JP2022501480A - Piezoelectric sensor containing electrospun poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHX) nanofibers - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a device comprising a PHA-based copolymer layer containing at least one of a ribbon or a polarized polymer composition of an electro-spun fiber of a polyhydroxyalkanoate-based copolymer obtained by the method according to claim 1 in the present specification. Is disclosed, the layers are configured to exhibit one or more of the piezoelectric, pyroelectric and strong dielectric effects, and the ribbon of electrospun fibers and each of the polarized polymer compositions are subjected to X-ray diffraction. Includes β-type PHA-based copolymers present in an amount of about 10% to about 99% as measured. The device may be configured for use as a sensor, actuator, nanomotor, or biobattery. [Selection diagram] FIG. 37.

Description

Cross-reference to related applications This application claims the priority of US Provisional Patent Application No. 62 / 734,360 filed on September 21, 2018, the entire disclosure of which is by reference in its entirety for all purposes. Incorporated herein.

Description of Federally Funded Research The invention was made with government support under grant number 1301765 awarded by Delaware NSF EPSCoR and grant number 1407255 awarded by the National Science Foundation through the DMR Polymers program. Government has certain rights in the present invention.

The term piezoelectricity, derived from the Greek word for pressure electricity, produces an electric charge when exposed to mechanical stress (direct piezoelectric effect) or expands or contracts when placed in an electric field (indirect / Inverse piezo effect), explaining the capabilities of a particular material. The piezoelectric effect is perceived as an electromechanical interaction between mechanical and electrical states in materials that do not have inversion symmetry. The nature of the piezoelectric effect is strongly related to the presence of the electric dipole moment in the solid, which can be induced by ions in the crystal lattice site surrounded by asymmetric charges or held directly by the molecular group. When a piezoelectric material is subjected to mechanical deformation, it exhibits a potential difference between its surfaces due to changes in the dipole moment. This potential difference or voltage circulates the charge in the circuit (current), thus producing electricity. Among all piezoelectric materials, piezoelectric polymers have been of particular interest over the last two decades due to their structural and dimensional flexibility, light weight, ease of processing, wide sensitive areas, and relatively low cost implementation. Has been done.

Polyhydroxyalkanoates (PHA) are a class of biodegradable and biocompatible aliphatic polyesters synthesized by various bacteria as intracellular carbon and energy storage materials. They are of scientific interest because of their promising environmental, electrical, pharmaceutical, and biomedical applications. Among PHA, poly (3-hydroxybutyrate) (PHB) homopolymers are the most common type and have been extensively studied over the last 30 years. However, due to the near-perfect stereoregularity, the PHB produced by bacteria has a very high crystallinity (> 60%) and a melting temperature range (about 180 ° C.) near its pyrolysis temperature. The difficult thermal properties of the material, as well as the stiffness and brittleness constraints, are major obstacles to most standard applications. Copolymerization with other low molecular weight monomer units, such as 3-hydroxyvalerate (3HV), has been attempted, but with relatively little success in improving properties. This surprising result is due to the fact that the 3HB and 3HV units are homogeneous diimages and the 3HV units are incorporated within the PHB crystal lattice. Recently, in order to substantially improve the properties of PHB, a small amount of hydroxyalkanoic acid monomer having a longer side chain, for example 3-hydroxyhexanoate (3HHx), has been copolymerized with 3HB units to obtain homogenous diimages. It was avoided and the rigidity and brittleness of the resulting copolymer was reduced. These medium chain length (mcl) branches act as molecular defects, disrupting the excessive regularity of the polymer chain and, as a result, reducing crystallinity and melting point (Tm). The resulting random copolymer, poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx), becomes flexible and flexible with increasing 3HHx content and is straightforward. It provides similar properties to chain low density polyethylene (LLDPE). Many properties of PHBHx, including chemical, thermal and mechanical properties, can be adjusted by varying the comonomer content.

PHB and poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyvalerate] (PHBV, PHB-based random copolymer) have two different types, α-type and β-type, depending on the processing conditions. It has been demonstrated that it can show crystalline polymorphism. Type α is the most common crystal structure of PHB or PHBV polymers obtained from typical crystallization processes such as melt or solution crystallization. This crystalline polymorph, the molecular chains take a left-handed 2 ₁ helical structure. The unit cell is orthorhombic, the space group is P2 ₁ 2 ₁ 2 _1- D ₂ , and the lattice constants are a = 0.576 nm, b = 1.320 nm, and c (fiber cycle) = 0.596 nm. be. Other polymorphs, β-type crystals, are recognized as strain-induced quasicrystal structures with highly elongated chains. In the β type, the chain has a twisted planar zigzag structure, which is almost completely elongated chain structure. The unit cell is also orthorhombic, and the lattice constants are a = 0.528 nm, b = 0.920 nm, and c (fiber cycle) = 0.470 nm. This metastable crystal structure, which can be reannealed to α-type at 130 ° C., was first observed in hot-stretched PHB thin films and then in cold-stretched PHBV thin films. In particular, this β-type is found in cold-stretched amorphous films, indicating that β-type formation does not require prior alignment of α-type crystals. Over the next decade, metastable β-types have been successfully produced in thin or melt-spun fibers of PHB and PHBV under different post-processing conditions, the membranes or fibers are highly pulled, but the draw ratios vary. Can be. It has been reported that the β-type remained relatively unchanged for several months at room temperature, suggesting that the crystal structure does not undergo secondary crystallization.

The β structure has long been accepted as a result of the orientation of free chains in the amorphous phase between α-type layered crystals. When a strong tensile force is applied, the tie molecules between the layered crystals are strongly elongated and oriented along the tensile direction. As long as the free chains have a planar zigzag structure, they fill to form a β structure. The formation of β-type crystal structures has a great effect on various properties of the material, including mechanical properties, biodegradability, and piezoelectric response.

Electrospinning is an effective and versatile technique that utilizes electrostatic forces to draw many different macromolecular solutions or melts to produce nanofibers (10 nm to 5 μm). Such fibers are found to be used in areas including composites, tissue engineering, energy storage and conversion, sensors, and filtration systems. Efforts have been made to elucidate the tensile forces generated by the strong electricity in the electrospinning process. The total draw ratio is estimated to reach 25,000. In addition, improved collectors such as rotating collectors (rotating drums, rotating discs, etc.) and gapped collectors (two charged metal rods or plates separated by an insulating gap) are added onto the fibers during fiber deposition. The tensile force of can be applied to finally obtain fibers that are macroscopically aligned along the winding direction or across the gap. It has been observed that these strong tensile forces induce the formation of metastable phases or polymorphs with very rapid evaporation of the solvent. Therefore, electrospinning has been investigated as a processing technique for inducing a metastable β structure in PHA nanofibers. The β structure was found in PHB nanofibers electrospun from a dilute polymer solution by conventional electrospinning techniques. Later, this metastable crystal structure was observed in electrospun PHBV fibers collected on a rotating drum. The presence of β-type in electrospinning PHBHx was described by Gong, Liang et al., “Discovery of β-form crystal structure in electrospinning poly [(R) -3-hydroxybutylate-co- (R) -3-hydoxy] From fiber mats to single fibers. ”, Macromolecules 48, 17 (2015): 6197-6205 and Gong, Liang et al.,“ Polymorphic Distribution in Independent Electrospinning- hydroxyhexanoate] (PHBHx) Nanofibers. ”, Macromolecules 50, 14 (2017): 551 to 5517.

In the art, it is still desired to develop applications for making, optimizing and using β-type PHBHx in various implementations.

As used herein, bio-based biodegradable poly (hydroxyalkanoate) copolymers (PHA) of 3-hydroxybutyrate and other 3-hydroxyalkanoate units with medium chain length (mcl) branching, such as poly [( For the first time, a sensor containing a successfully produced material containing [R) -3-hydroxybutyrate-co- (R) -3-hydroxyalkanoate] was disclosed, which material has a surprisingly high content of β. It has a type crystal structure and the extended chain has a planar zigzag structure verified by wide-angle X-ray diffraction (WAXD) analysis. In addition, these materials, including high content β-type crystals, exhibit unexpectedly high levels of piezoelectricity, pyroelectricity and ferroelectricity. Therefore, these materials can be used to manufacture devices such as actuators, sensors and the like. In addition, materials containing mcl-branched PHA have substantial advantages over conventional PHA due to the superior properties that make them more suitable for many end applications and the much easier processability. This β-type crystal structure was produced under selected very specific conditions according to various embodiments according to the present invention. Process methods include high-speed rotating disk fiber collection with air gaps or sharp edges to receive electrodes, stretching of extremely thin films, shearing of thin films between two glass slides, electrospinning of thin films, high pressure annealing, in thermoreversible gels. Can include highly aligned electrospun PHA nanofiber manufacturing techniques, including nanofibril formation and high pressure or high temperature treatment.

Also disclosed herein are strain-induced, semi-stable β-type crystals, in which the elongated chains were collected over the air gaps on the aluminum foil and on the tapered edges of the fast rotating disk, poly [(R). The macroscopically aligned electrospun nanofibers of -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx) have a planar zigzag structure. The presence of β-type crystal structure in the fiber mat was confirmed by wide-angle X-ray diffraction (WAXD) and Fourier transform infrared spectroscopy (FTIR). In addition, selected area diffraction (SAED) and AFM-IR were used to investigate the morphological and structural details of individual electrospun nanofibers. The results of SAED confirmed the significant effect of the collection method on the crystal structure and the alignment level of the molecular chains in the crystal. The AFM-IR spectrum of a single nanofiber was in good agreement with the conventional FTIR spectrum, but the finer features in the AFM-IR spectrum were more clearly and better resolved. Based on the experimental results, a new mechanism for the formation of β-type crystal structure in electrospun PHBHx nanofibers is proposed.

In one embodiment, the β form of PHBHx also exhibits pyroelectric and ferroelectric properties.

In one aspect of the invention, a polymer blend exhibiting at least one of piezoelectric, pyroelectric or ferroelectric properties is provided, which is intended for use in devices such as sensors and actuators. In one embodiment, the polymer blend is a blend of one or more PHBHx copolymers, each of which has a different comonomer content. For example, the polymer blend may include a blend of PHBHx3.9 and PHBHx13.0 as an example of a piezoelectric blend. In another embodiment, the polymer blend is a polymer blend of PHBHx and a polymer having a different composition. For example, polymer blends may include, but are not limited to, polymer blends of PHBHx with poly (vinylidene fluoride) (PVF2) and copolymers thereof, and / or nylon 5, nylon 7, nylon 9, nylon 11, and the like. ..

Also disclosed herein are PHAs with longer side chains such as pentyl, heptyl, etc., including, but not limited to, Nodax ™ class PHA.

In general, the PHA copolymer according to the present invention is preferably in the range of about 10% to about 99% as measured by X-ray diffraction, more preferably from about 20% to about 80% based on X-ray diffraction analysis. More preferably, it has a β type present in the range of about 30% to about 70%.

In one aspect, one of the following methods can be used to make a PHA copolymer-based piezoelectric material.
-Fibers produced by the electrospinning process (currently known to be the best method)
-Film or sheet subjected to high shearing, eg calender rolling-Film or sheet made by crystallization of melt under shearing or pressure-Gel dried under shearing pressure-Clearing by solvent crystallization Gels frozen to induce-articles processed under shear or pressure-articles made from melts in an electric field (polarization)
-Other processing methods for increasing piezoelectric PHBHx include extrusion (eg, in a twin-screw extruder), injection molding, and conventional fiber spinning.

As used herein, single electrospun polymer nanofibers containing an inhomogeneous spatial distribution of crystalline polymorphs are disclosed. Are two crystalline polymorphs of PHBHx, 2 ₁ thermodynamically stable α-type consisting of a chain having a helical structure, and consists of a chain having a planar zigzag structure metastable β type core rich in α-type and It is spatially distributed as a core shell structure composed of β-type rich shells. Furthermore, the thickness of the shell was found to be independent of the fiber diameter. The determination of the crystal polymorphic distribution in individual nanofibers was made possible by using a technique that combines atomic force microscopy (AFM) and infrared spectroscopy (IR), which is a nanoscale space. It provides resolution and crystal phase specificity at the same time. Based on the experimental results, a possible generation mechanism of this polymorphic inhomogeneous core-shell structure is proposed. The involvement of this core-shell model in fiber properties is also discussed.

Disclosed herein are piezoelectric electrospun PHBHx nanofibers having a crystalline region exhibiting a β-type crystal structure associated with a chain of highly oriented planar zigzag structure. The piezoelectric properties of the aligned PHBHx nanofibers were investigated as a function of the fluctuating crystal structure. The results showed a strong correlation between the piezoelectric response of the fiber and the presence of the β-type crystal structure. Although the mechanism by which the piezoelectric response of the fiber is generated has been proposed, the sensitivity of the piezoelectric PHBHx nanofiber to pressure has also been quantified.

In the present specification, a method for forming a β structure in a PHBHx film by stress-induced β crystallization is disclosed, and stress is applied by mechanically pulling the film.

In the present specification, a method for forming a β structure in a PHBHx film by a novel method of mechanical tension following room temperature isothermal crystallization is disclosed. It was confirmed that α crystallites must be present before the formation of the β structure. The ideal conditions for the initial formation of the β structure corresponded to 28 minutes of isothermal crystallization. Furthermore, it was shown that this β-structure is reversible and that the tensile process is different in the initial β-forming and re-tensile processes. The β structure can also be annealed to the α structure again at a temperature as low as 48 ° C.

In one embodiment, a method for preparing a polarized polymer composition.
a) A step of providing a layer of a polarizable polymer composition, wherein the polarizable polymer composition comprises a polyhydroxyalkanoate-based copolymer;
b) The process of directionally perturbing the layer to induce polarization;
c) Optionally, a step of polarizing the polymer composition of the directionally perturbed layer by applying a high electric field with a strength lower than the strength that results in substantial dielectric breakdown of one or more polymers;
d) Optionally, the step of annealing the polarized polymer composition of the layer at a temperature lower than the melting temperature of the crystals of the polarized polymer composition, whereby the polarization is maintained up to the crystal melting point of the polar crystals of the polymer composition. A method is provided in which the polarized polymer composition comprises a β-type PHA-based copolymer present in an amount of about 10% to about 99% as measured by X-ray diffraction.

In one embodiment of the method, the step of providing the layer comprises electrospinning a ribbon of fibers from a solution of the polyhydroxyalkanoate-based copolymer in one or more solvents, of the ribbon of electrospun fibers. Each fiber contains a β-type shell and an α-type core.

In one embodiment of the method, the step of providing the layer comprises forming a layer from a solution of the polyhydroxyalkanoate-based copolymer in one or more solvents.

In another embodiment of the method, the step of providing the layer comprises forming the layer from the melt composition of the polyhydroxy alkanoate-based copolymer, and the step of directionally perturbing the layer is the calendering of the layer after quenching. Includes shearing or cold stretching.

In one embodiment of the method, the step of providing the layer comprises forming a membrane from the gel composition of the polyhydroxyalkanoate-based copolymer, and the step of directionally perturbing the membrane dried the gel under shear pressure. It involves allowing or freezing the gel to induce shear by crystallization of the solvent.

In one embodiment of the method, the step of providing the layer comprises forming an electrospun fiber mat from a solution of the polyhydroxyalkanoate-based copolymer in one or more solvents.

In one embodiment, the polyhydroxyalkanoate-based copolymer is a hydroxybutyrate unit, a hydroxyhexanoate unit, a vinyl unit, a vinylidene unit, an ethylene unit, an acrylate unit, a methacrylate unit, a nylon unit, a carbonate unit, an acrylonitrile unit, a cellulose unit. , Pendant fluoro, chloro, amide, acrylate and ester other than the methacrylate unit ester, cyanide, nitrile other than the acrylonitrile unit, or at least one selected from the group consisting of units having an ether group, protein units, and combinations thereof. Contains one monomer unit.

In another embodiment, the polarizable polymer composition is polyvinyl chloride, polymethylacrylate, polymethylmethacrylate, poly (vinylidene cyanide / vinyl acetate) copolymer, vinylidene cyanide / vinyl benzoate copolymer, vinylidene cyanide /. Isobutylene copolymer, vinylidene cyanide / methyl methacrylate copolymer, vinylidene fluoride copolymer, polyvinyl chloride, polyacrylonitrile, polycarbonate, cellulose, protein, synthetic polyester and ether of cellulose, poly (γ-methyl-L-glutamate), vinylidene copolymer, Further comprising one or more polymers selected from the group consisting of Nylon-3, Nylon-5, Nylon-7, Nylon-9, Nylon-11 and blends thereof.

In one aspect of the method, the layer comprises a polarizable polymer composition of at least two layers formed from the multiphase composition of said polymer composition. In one embodiment, at least two layers are coextruded and in contact with each other.

In one embodiment of the method, the step of polarization the polymer composition of an optionally directionally perturbed layer is up to about 20 ° C. to about 120 ° C. using an electric field of at least 1 MV / cm. It takes place for 5 hours. In another embodiment, the layer is annealed at a temperature in the range of about 125 ° C to about 150 ° C for at least 1 hour.

In one aspect of the invention, there is provided a device comprising at least one electrospun fiber mat or polarized polymer composition of a polyhydroxyalkanoate-based copolymer obtained by the method of claim 1, wherein the device has a piezoelectric effect. , One or more of the pyroelectric and ferroelectric effects.

In one embodiment of the device, the device further comprises two or more layers of the polarized polymer composition, the two or more layers being in the form of a ribbon of fibers laminated to each other.

In one aspect, the device is a sensor configured to generate a potential difference or voltage in response to changes in the dimensions of the layers of the polarized polymer composition. In one embodiment, the dimensional change is one or more of the following properties on the surface of the sensor layer of the polarized polymer composition: humidity, temperature, salt, nutrient attachment or injection, and metalloid attachment: Caused by.

In another aspect of the device, the sensor comprises multiple sensor surfaces and / or interfaces, each surface / interface having the following characteristics: humidity, temperature, salt, nutrient attachment or injection, and metalloid attachment. Independently configured to monitor one of them.

In one aspect, the device is an actuator configured to expand or contract in response to an application of charge across layers of the polarized polymer composition.

In one aspect, a nanomotor comprising the one or more piezoelectric actuators described herein is provided.

In another embodiment, it is configured to be placed or implanted near the vasculature components of an animal or human patient, and the heartbeat of the animal or human patient produces and monitors dimensional changes in the PHA-based copolymer layer. Alternatively, a device comprising one or more actuators disclosed above disclosed herein is provided that is configured to generate a potential difference or voltage for manipulating the treatment device. In one embodiment, the treatment or monitoring device includes a pacemaker, insulin pump, in-situ glucose monitor, or blood pressure monitor. In another embodiment, the vasculature component comprises a vein or artery that beats in a rhythm corresponding to the patient's heartbeat.

In one embodiment, a method comprising providing a PHA-based copolymer sensor layer configured to monitor one of the following properties: humidity, temperature, salt content, nutrient attachment or injection, and semi-metal adhesion: Provided, the PHA-based copolymer sensor layer comprises at least one ribbon or polarized polymer composition of electrospun fibers of the polyhydroxyalkanoate-based copolymer obtained by the methods disclosed above. The method comprises exposing the PHA-based copolymer sensor layer to one or more of the properties such that the PHA-based copolymer sensor layer undergoes a dimensional change caused by one or more of the properties. It further comprises the step of detecting a potential difference or voltage in response to changes in the dimensions of the PHA-based copolymer sensor layer.

SEM images of electrospun PHBHx fibers collected on (a) aluminum foil outside the air gap, (b) over the air gap, and (c) on a rotating disk. It is a figure which shows the polarization FTIR spectrum of the Al foil random fiber (a), the air gap alignment fiber (b), and the rotating disk alignment fiber (c). The incident IR beam is polarized in two directions perpendicular to each other, perpendicular (red) and parallel (black) to the aligned fiber axis. It is a figure which shows the WAXD profile of the electrospun PHBHx fiber obtained using different collectors. It is a figure which shows the FT-IR spectrum of the random fiber outside the gap, the alignment fiber in a gap, and the alignment fiber in a rotating disk in (a) a CDO expansion / contraction region and (b) a C—O—C expansion / contraction region. It is a figure which shows the bright-field TEM image of the single field-spun PHBHx nanofiber from different collectors (a-c), and their corresponding SAED pattern (a'-c'). (A) AFM images of single electrospun PHBHx fibers collected on aluminum foil (473 nm) and (b) on tapered edges (324 nm) of a rotating disk. (C) IR spectra of the fibers shown in (a) and (b). Red dots on the two individual fibers indicate the location of the AFM tip. It is a figure which shows two possible generation mechanism of α-type crystal structure and β-type crystal structure during electric field spinning and collection. Path 1 shows the simultaneous formation of α-crystal morphology and β-crystal morphology during nanofiber collection. Path 2 shows that the α crystal morphology is generated after the formation of β type. The α crystal morphology results from a planar zigzag chain (β type) relaxed by the presence of residual solvent in the core. In this figure, broken line in the fiber shows a 2 ₁ helix backbone chain of α Aiuchi, straight line shows a plan zigzag backbone chain of β Aiuchi, a random curves were free amorphous Aiuchi Shows a chain. The cyan color in the figure represents a solvent. It is a figure which shows the AFM height image (a, a'), the IR peak image (b, b'), and the AFM-IR spectrum (c, c') of two single electric field spinning PHBHx fibers. The black spots in (b) and (b') indicate the position of the AFM tip when collecting the AFM-IR spectrum. It is a figure which shows the original (a, a') and contrast inversion (b, b') SAED patterns of two individual field-spun PHBHx nanofibers of 251 nm and 417 nm, respectively. The inset is a brightfield TEM image of the two fibers. The intensity profiles of (a) and (a') along the equator are plotted in (c) and (c'), respectively. FIG. 5 shows a comparison of AFM-IR spectra collected from different positions on two individual electrospun PHBHx nanofibers. It is a figure which shows the AFM height image (a, a') and IR peak image ratio (b, b') of the cross section of two individual electric field spinning PHBHx nanofibers. It is a figure which shows the possible generation mechanism for the formation of the core-shell structure of the electrospun PHBHx nanofibers collected by a rotating disk. The wavy line in the core of the final solidified fibers showed 2 ₁ helix backbone chain in the α crystalline form, straight lines in the shell of the fiber and the final solidification fiber immediately after spinning, the planar zigzag chains in β-crystal form It shows the skeleton and the random curves show the free chains in the amorphous region. The blue color in this figure represents the solvent. The thickness of the shell is exaggerated for display purposes. It is a figure which shows the schematic (a) of the piezoelectric cantilever and the plot (b) of the displacement of a probe with respect to the applied voltage. It is a schematic diagram of the piezoelectric response test equipment. FIG. 3 shows SEM images and WAXD patterns of macroscopically aligned PHBHx nanofibers before (a, a') and after (b, b') annealing at 130 ° C. for 24 hours. It is a figure which shows the induced voltage vs. time output of the rotating disk alignment PHBHx nanofiber before annealing (black signal) and after annealing (red signal). It is a figure which shows the electric dipole in (a) 2 ₁ spiral PHB chain (α type) and (b) plane zigzag PHB chain (β type). It is a figure which shows the charge transfer during a press-hold-release cycle. The black line represents the actual voltage signal. The dark green line represents the free charge transport from the external circuit to the fiber. The pink line represents the net charge on the ribbon of the fiber. (A) It is a schematic diagram of the analysis model of the response of the ribbon of the fiber of PHBHx under the applied pressure. The blue letters p and the red letters A indicate the applied pressure and the contact area between the probe and the ribbon of the fiber, respectively. (B) Experimental (black) and linear fitting (red) pressure response curves. WaxD profiles of untreated annealed 3.9 (a) and 13 (b) mol% 3HHx PHBHx films, as well as (c) 13 mol% 3HHx PHBHx films oriented and pulled perpendicular to the tensile direction with an X-ray beam. It is a figure which shows. FIG. 5 shows an IR spectrum of an untreated 13 mol% 3HHx PHBHx polymer membrane as a function of increasing strain. It is a figure which shows the Raman spectrum of the untreated 13 mol% 3HHx PHBHx polymer membrane before and after the strain. It is a figure which shows the Raman spectrum of the untreated 13 mol% 3HHx PHBHx polymer membrane before and after the strain, centering on a carbonyl region. It is a figure which shows the pulled membrane which shows the α region which was not pulled and the β region which was pulled. It is a figure which shows the XRD of the amorphous non-pulled α region and the pulled β region of the same film. It is a figure which shows the XRD spectrum of the membrane after being pulled and after being released. It is a figure which shows the XRD spectrum of the membrane after being pulled, after being released, and after the tension is restored. It is a figure which shows XRD as a function of distortion. It is a figure which shows the XRD spectrum of the stretched membrane which was annealed at 48 degreeC for 2 hours. It is a figure which shows the XRD spectrum of the pulled film after isothermal crystallization at room temperature. It is a figure which shows the Raman spectrum of the pulled film after isothermal crystallization at room temperature. It is a figure which shows the thermoreversible sol-gel transition in PHBHx CHCl _{3 / DMF solution.} It is a figure which shows the WAXD profile of (a) raw PHBHx powder and (b) lyophilized gel. It is a figure which shows the WAXD profile of the gel film before (a) and after (b) of tension. It is a figure which shows the scanning electron microscope image of PHBHx3.9 mol% electrospun nanofibers collected on a rotating wheel (3500 rpm) showing the β crystal form by wide-angle X-ray scattering. FIG. 5 shows aligned PHBHx3.9 mol% nanofibers (see red ellipse) of FIG. 35 as part of an electrical circuit deformed by a pliers handle to generate a voltage to light an LED bulb. FIG. 3 is a schematic diagram of a sensor / actuator device according to an exemplary embodiment of the present invention. It is a schematic diagram which shows the flowchart which concerns on the exemplary method which uses the exemplary sensor embodiment of this invention. It is a schematic diagram which shows the flowchart which concerns on the exemplary method which uses the exemplary sensor embodiment of this invention in a biological use.

As used herein, "PHA" means the polyhydroxyalkanoate according to the invention. As used herein, "PHB" means poly (3-hydroxybutyrate), which is a homopolymer. As used herein, "PHBV" means the copolymer poly (3-hydroxybutyrate-co-3-hydroxyvalerate). As used herein, "PHBHx" means the copolymer poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate], where x is the copolymer. It represents the amount of hydroxyhexanoate (Hx), which is a copolymer present therein, as mol%.

In one embodiment, the PHBHx copolymer is produced by chiral ring-opening polymerization of optically pure 3HB and 3HHx [R] structural comonomer. In another embodiment, the PHBHx copolymer is a biologically produced biobased isotactic PHBHx.

Β-type crystal structure of electrospun poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx) nanofibers: from fiber mats to single fibers. Disclosed in β crystal polymorphs in macroscopically aligned electrospun PHBHx nanofibers obtained using two improved collectors, an aluminum foil with a rectangular air gap and a rotating disk with tapered edges. Will be done. The fiber morphology, crystal structure, and chain structure were characterized by scanning electron microscopy (SEM), wide-angle X-ray diffraction (WAXD), and transmission-type Fourier transform infrared spectroscopy (FTIR) on the fiber mat. In addition, selected area diffraction (SAED) and AFM-IR were used to investigate the structure and orientation within PHBHx nanofibers on a single fiber scale. Example 1 later disclosed herein is β-type in electrospun poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx) nanofibers. Exemplary methods of forming crystal structures are provided, including experimental procedures and results.

Piezoelectric Broadhurst and Davis (Broadhurst, MG and G.) in electrospun poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx) nanofibers. According to T. Davis, "Piezo-and hydroxy products." Electrics. Springer, Berlin, Heidelberg, 1980. 285-319), all piezoelectric polymers have the following four important elements: (1). Permanent molecular dipole generation, (2) ability to align molecular dipoles, (3) ability to maintain this dipole alignment once achieved, and (4) large strain when mechanical stress is applied. The ability of the material to produce. Electrospun PHBHx nanofibers (1) for O = C-CH ₂ dipoles are present in the material, (2) the molecular chains are oriented along the fast and extensive tensile by fiber axis in the electrospinning, Therefore, because the molecular dipoles are aligned, (3) rapid solvent evaporation during electrospinning facilitates the preservation of the dipole alignment through fiber solidification, and (4) the flexible fibers withstand large deformations. Therefore, it meets all the above criteria.

In one aspect, the piezoelectric response of electrospun PHBHx nanofibers depending on the crystal structure present in the fiber is disclosed herein. Macroscopically aligned PHBHx nanofibers containing a metastable β-type crystal structure were produced using a high speed rotating disk as a collector. Control samples were made by annealing these fibers at 130 ° C. for 24 hours so that the annealed fibers contained only α crystal morphology. A piezoelectric cantilever was used to characterize the piezoelectric properties of the aligned fibers both before and after annealing. The mechanism by which the piezoelectric response of the fiber is generated has been proposed, but the sensitivity of the piezoelectric PHBHx nanofiber to pressure has been quantified. Example 3 later disclosed herein provides an exemplary method for measuring the piezoelectric effect in electrospun PHBHx nanofibers.

In one embodiment, the β form of PHBHx also exhibits pyroelectric and ferroelectric properties.

In one embodiment, a hydroxybutyrate unit, a hydroxyhexanoate unit, a vinyl unit, a vinylidene unit, an ethylene unit, an acrylate unit, Polymer units, nylon units, carbonate units, acrylonitrile units, cellulose units, pendant fluoros, chloros, amides, acrylates and methacrylate units other than esters, cyanides, nitriles other than acrylonitrile units, or units with ether groups, protein units , As well as polyhydroxyalkanoate (PHA) copolymers containing polarizable polymer units selected from the group consisting of combinations thereof.

In another embodiment, poly (hydroxybutyrate / hydroxyhexanoate), poly (vinylidene fluoride / trifluoroethylene), poly (vinylidene fluoride / tetrafluoroethylene), poly (vinylidene fluoride / vinyl trifluoride). , Polyethylene copolymers selected from the group consisting of poly (vinylidene fluoride / vinyl chloride) and poly (vinylidene fluoride / methylmethacrylate) are provided.

Copolymers can be up to 1%, 2%, 3%, 5%, 10%, 15%, 20%, 25%, 30%, 35%, or 2-30% or 3-25% or 5-20%. It may be present in the polyhydroxyalkanoate copolymer in any suitable molar amount, including within the range. In one embodiment, the comonomer unit hydroxyhexanoate is any suitable mole comprising up to 1%, 2%, 3%, 5%, 10%, 15%, 20%, 25%, 30%, 35%. Present in poly (hydroxybutyrate / hydroxyhexanoate) in quantity.

In yet another embodiment, polyvinyl chloride, polymethylacrylate, polymethylmethacrylate, poly (vinylidene cyanide / vinyl acetate) copolymer, vinylidene cyanide / vinyl benzoate copolymer, vinylidene cyanide / isobutylene copolymer, vinylidene cyanide / methyl. Polarization containing one or more polymers selected from the group consisting of methacrylate copolymers, polyvinyl chloride, polyacrylonitrile, polycarbonate, cellulose, proteins, synthetic polyester and ethers of cellulose, and poly (γ-methyl-L-glutamate). Possible polymer compositions are provided.

In one embodiment, the polarizable polymer composition comprises one or more polymers selected from the group consisting of poly (vinylidene fluoride) and vinylidene fluoride copolymers. In another embodiment, the polymerizable polymer composition is poly (hydroxybutyrate / hydroxyhexanoate), poly (vinylidene fluoride / trifluoroethylene), poly (vinylidene fluoride / tetrafluoroethylene), poly (vinylidene fluoride / tetrafluoroethylene). Includes one or more copolymers selected from the group consisting of vinylidene fluoride / vinyl trifluoride), poly (vinylidene fluoride / vinyl chloride) and poly (vinylidene fluoride / methylmethacrylate). In yet another embodiment, the polarizable polymer compositions are soluble ceramic materials, PHBHx, poly (vinylidene fluoride), vinylidene copolymers, nylon-3, nylon-5, nylon-7, nylon-9 and nylon-11. It also comprises one or more polymers selected from the group consisting of blends thereof.

In one aspect of the invention, the polarizable polymer composition is a polymer blend exhibiting at least one of piezoelectric, pyroelectric or ferroelectric properties, the polymer blend being used in devices such as sensors and actuators, for example. Is for. In one embodiment, the polymer blend is a blend of one or more PHBHx copolymers, each of which has a different comonomer content. For example, the polymer blend may include a blend of PHBHx3.9 and PHBHx13.0 as an example of a piezoelectric blend. In another embodiment, the polymer blend is a polymer blend of PHBHx and a polymer having a different composition. For example, polymer blends may include, but are not limited to, polymer blends of PHBHx with poly (vinylidene fluoride) (PVF2) and copolymers thereof, and / or nylon 5, nylon 7, nylon 9, nylon 11, and the like. ..

In another aspect of the invention, the polarizable polymer composition is PHA with longer side chains such as pentyl, heptyl and the like. In one embodiment, the polymerizable polymer composition is a polymer blend of PHBHx and one or more PHA with longer side chains. In another embodiment, the polarizable polymer composition is a copolymer of PHBHx with one or more PHA with longer side chains. Exemplary PHA with longer side chains are from, but not limited to, PHBOs with pendant side groups, PHBDs with heptyl side chains, and similar ones with side groups such as the C15 side chain, Danimer Scientific. Includes the available Nodax ™ class medium chain long branched polyhydroxyalkanoates, mcl-PHA. For a list above PHBHx, see US5,602,227 and RE36,584, which are incorporated herein by reference.

In one embodiment, the polarizable polymer composition is a blend of PHBHx, poly (vinylidene fluoride) and vinylidene fluoride-vinyl fluoride (80/20) copolymer. In another embodiment, the polarizable polymer composition is a blend of 50:50 by weight PHBHx with a vinylidene fluoride-vinyl fluoride (80/20) copolymer. In another embodiment, the polarizable polymer composition is a blend of PHBHx by weight 50:50 and polyvinylidene fluoride.

Β-type in blends of PHA copolymers and PHA copolymers In general, blends of PHA copolymers and PHA copolymers according to the present invention are preferably in the range of about 10% to about 99% as measured by X-ray diffraction, more preferably. It has a β-type present in the range of about 20% to about 80%, more preferably more preferably about 30% to about 70% based on X-ray diffraction analysis.

Processes for Making PHA Copolymer-Based Piezoelectric Materials Under Directional Perturbation Any suitable method can be used to make PHA copolymer-based piezoelectric materials, including but not limited to:
-Fibers produced by the electrospinning process (currently known to be the best method).
-Membranes or sheets subjected to high shear, eg calendar rolling. Example 5, later disclosed herein, provides an exemplary method for the formation of β-type PHBHx copolymer-based membranes by mechanical tension.
-Membrane or sheet made by crystallizing the melt under shear or pressure. Example 4, later disclosed herein, provides an exemplary method of stress-induced β-crystallization of PHBHx copolymer-based membranes.
-A gel that dries under shear pressure. Example 6 later disclosed herein provides an exemplary method for the formation of β-type PHBHx copolymer-based gel membranes by mechanical tension.
-A gel that is frozen to induce shear by crystallization of the solvent.
-A molded article that is processed under shear or pressure.
-Articles made from melt in an electric field.
-Other processing methods for increasing piezoelectric PHBHx include extrusion (eg, in a twin-screw extruder), injection molding, and conventional fiber spinning.

In one embodiment, a method for preparing a polarized polymer composition.
a) A step of providing a layer of a polarizable polymer composition, wherein the polarizable polymer composition comprises a polyhydroxyalkanoate-based copolymer;
b) The process of directionally perturbing the layer to induce polarization;
c) Optionally, a step of polarizing the polymer composition of the directionally perturbed layer by applying a high electric field with a strength lower than the strength that results in substantial dielectric breakdown of one or more polymers;
d) Optionally, a step of annealing the polarized polymer composition of the layer at a temperature lower than the melting temperature of the crystals of the polarized polymer composition, whereby the polarization is maintained to the crystal melting point of the polar crystals of the polymer composition. Process and
A method is provided in which the polarized polymer composition comprises a β-type PHA-based copolymer present in an amount of about 10% to about 99% as measured by X-ray diffraction.

In one embodiment of the method, the step of providing the layer comprises electrospinning a ribbon of fibers from a solution of the polyhydroxyalkanoate (PHA) based copolymer in one or more solvents, the electrospun fibers. Each fiber of the ribbon contains a β-type shell and an α-type core. In one embodiment, up to 1 wt% or up to 2 wt% or up to 5 wt% PHA in a dilute solution, such as a suitable solvent such as 1,1,1,3,3,3-hexafluoro-2-propanol (HFIP). The base copolymer can be used for electrospinning using experimental configurations as disclosed herein in Example 1.

In one embodiment of the method, the step of providing the layer comprises forming a layer from a solution of the polyhydroxyalkanoate-based copolymer in one or more solvents. The polyhydroxyalkanoate-based copolymer may be present in an amount of 1-99% or 5-90% or 10-80% based on the total weight of the solution. Any suitable solvent including, but not limited to, chloroform, toluene, acetone and the like can be used. The layer can be formed using any suitable method such as casting.

In another embodiment of the method, the step of providing the layer comprises forming the layer from the melt composition of the polyhydroxy alkanoate-based copolymer, and the step of directionally perturbing the layer is the calendering of the layer after quenching. Includes shearing or cold stretching.

In one embodiment of the method, the step of providing the layer comprises forming a membrane from the gel composition of the polyhydroxyalkanoate-based copolymer, and the step of directionally perturbing the membrane dried the gel under shear pressure. It involves allowing or freezing the gel to induce shear by crystallization of the solvent. The gel composition of the polyhydroxy alkanoate-based copolymer can be formed by any suitable method as disclosed later in Example 6 herein.

In one embodiment of the method, the polyhydroxyalkanoate-based copolymer is a hydroxybutyrate unit, a hydroxyhexanoate unit, a vinyl unit, a vinylidene unit, an ethylene unit, an acrylate unit, a methacrylate unit, a nylon unit, a carbonate unit, an acrylonitrile unit, Selected from the group consisting of esters other than esters of cellulose units, pendant fluoros, chloros, amides, acrylates and methacrylate units, cyanides, nitriles other than acrylonitrile units, or units having ether groups, protein units, and combinations thereof. Contains at least one monomer unit.

In another embodiment of the method, the polarizable polymer composition is polyvinyl chloride, polymethylacrylate, polymethylmethacrylate, poly (vinylidene cyanide / vinyl acetate) copolymer, vinylidene cyanide / vinyl benzoate copolymer, cyanated. Vinylidene / Isobutylene Copolymer, Vinylidene Cyanide / Methyl Methacrylate Copolymer, Vinylidene Fluoride Copolymer, Vinyl Fluoride, Polyacrylonitrile, Polycarbonate, Cellulose, Protein, Synthetic Polyester and Cellular Ether, Poly (γ-Methyl-L-Glutamate), Vinylidene Further comprising one or more polymers selected from the group consisting of copolymers, nylon-3, nylon-5, nylon-7, nylon-9, nylon-11 and blends thereof.

In one aspect of the method, the layer comprises a polarizable polymer composition of at least two layers formed from the polymorphic composition of said polymer composition. In one embodiment, at least two layers are coextruded and in contact with each other.

In one embodiment of the method, the step of polarization the polymer composition of an optionally directionally perturbed layer is up to about 20 ° C. to about 120 ° C. using an electric field of at least 1 MV / cm. It takes place for 5 hours. In one embodiment, the layer is annealed at a temperature in the range of about 125 ° C to about 150 ° C for at least 1 hour.

In one embodiment, the polarized polymer composition is a polarized poly (hydroxybutyrate-co-hydroxyhexanoate) (PHBHx) copolymer, and the polarization is essentially stable to the extent to which the polar crystals melt. .. In another embodiment, the polarized polymer composition is a polarized poly (hydroxybutyrate) / poly (vinylidene fluoride) copolymer.

In one embodiment, the polarized polymer composition comprises a polarized polyhydroxyalkanoate (PHA) based copolymer and poly (hydroxybutyrate / hydroxyhexanoate), poly (vinylidene fluoride / trifluoroethylene), poly (fluoride). In a blend with a copolymer selected from the group consisting of vinylidene / tetrafluoroethylene), poly (vinylidene fluoride / vinyl trifluoride), poly (vinylidene fluoride / vinyl chloride) and poly (vinylidene fluoride / methylmethacrylate). be. In another embodiment, the polarized polymer composition comprises a group consisting of a PHBHx copolymer and a soluble ceramic material, polyvinylidene fluoride, vinylidene copolymer, nylon-3, nylon-5, nylon-7, and nylon-11. A blend with one or more selected ingredients. In one embodiment, the polarized polymer composition comprises 50:50 by weight Nylon-7 and Nylon-11, respectively, as well as a polarized PHA-based copolymer. The polarized PHA-based copolymer may be present in any suitable amount in the blend, eg, in an amount of 1-99% or 5-90% or 10-80% relative to the total weight of the blend composition.

In one embodiment, the polarized polymer composition is a blend of PHBHx, poly (vinylidene fluoride) and vinylidene fluoride-vinyl fluoride (80/20) copolymer. In another embodiment, the polarized polymer composition comprises a blend of PHBHx at 50:50 by weight with a vinylidene fluoride-vinyl fluoride (80/20) copolymer. In another embodiment, the polarized polymer composition comprises a blend of PHBHx by weight of 50:50 with poly (vinylidene fluoride).

In one aspect of the invention, there is provided a device comprising at least one electrospun fiber mat or polarized polymer composition of a polyhydroxyalkanoate-based copolymer obtained by the method of claim 1, wherein the device has a piezoelectric effect. , One or more of the pyroelectric and ferroelectric effects. In one embodiment, the layer of the polarized polymer composition is a self-supporting sheet of the polarized polymer composition. In another embodiment, the layer of the polarized polymer composition is a non-self-supporting layer of said composition disposed on a support substrate.

In one embodiment of the device, the device further comprises two or more layers of the polarized polymer composition, the two or more layers being in the form of a ribbon of fibers laminated to each other.

In one aspect, the device is a sensor configured to generate a voltage in response to changes in the dimensions of the layers of the polarized polymer composition.

In another embodiment, the device is caused by one or more changes in the following properties on the surface of the sensor layer of the polarized polymer composition: humidity, temperature, salt content, nutrient attachment or injection, and metalloid attachment: A general purpose sensor configured to generate a potential difference or voltage in response to changes in dimensions.

In another aspect of the device, the sensor comprises multiple sensor surfaces and / or interfaces, each surface / interface having the following characteristics: humidity, temperature, salt, nutrient attachment or injection, and metalloid attachment. Independently configured to monitor one of them.

In one aspect, the device is an actuator configured to expand or contract in response to an application of charge across layers of the polarized polymer composition.

In one aspect, a nanomotor comprising the one or more piezoelectric actuators described herein is provided.

Further, in the present specification, an example of a method for producing a so-called polarized polymer article (electret) will be discussed later.

As used herein, a method is disclosed in which a highly polarized material can be produced, wherein the material has no or substantially no mechanically induced orientation, and the polarization is that of the polarized material. It is inherently stable up to the near crystal melting temperature range (or glass transition temperature) or, in the case of non-crystalline polarization materials, up to the near softening temperature range or the glass transition temperature range of the polarization material. The method comprises dissolving the polarized material in a solvent or solvent for the material. A solvent is selected that is adapted to the polarization of the material and can be removed to the desired extent during or after polarization during evaporation. The temperature used can be a temperature at which polarization is effective, usually a high temperature at which substantial dielectric breakdown does not occur. The DC electric field used for polarization can be selected to provide the desired polarization. Example 7, later disclosed herein, provides an exemplary method of making a polarized PHA copolymer-based article.

Also provided herein is a polarized product, which has no or substantially no mechanically induced orientation, up to near the crystalline melting point of the material, or. In the case of amorphous materials, they are inherently stable up to the near softening temperature range or glass transition temperature range of the material. Currently preferred materials are poly (hydroxyalkanoate) copolymers and, as a particular example, poly (hydroxybutyrate) copolymers.

In one embodiment of the invention, a polarized polymer composition characterized as follows is provided.
-Without or substantially no mechanically induced molecular orientation;
-Having a polarization that is inherently stable up to the near crystal melting temperature range of the polar crystals of the material, or to the near softening temperature range or the glass transition temperature range of the material if the material is anisotropic; and- Polarized It has isotropic or substantially isotropic mechanical and electromechanical properties in a plane perpendicular to the direction of the polarized electric field; the polarized material has a composition comprising essentially all or all of the material. It is a thing.

In one embodiment, a single electrospun fiber with PHBH x 3.9 mol% produced a piezoelectric response with a peak-to-peak response of 230 millivolts.

Electro-spun poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx) nanofiber introduction part For self-assembly by electric charge and its interaction with the applied electric field Based electrospinning is an efficient and versatile technique for producing ultrafine fibers with diameters up to tens of nanometers. Due to the strong tensile force and rapid solvent evaporation reaction rate associated with the electrospinning process, electrospun fibers may have different crystallization behavior compared to bulk materials. This can result in the formation of metastable phases or polymorphs. For some polymer materials, more than one crystalline polymorph can be found in electrospun nanofibers, the number of each polymorph can often be controlled by changing electrospinning conditions. For example, in an electrospun nanofiber of bio-based poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx) collected in an improved collector, two crystals. Coexistence of polymorphs has been observed. Further, thermodynamically stable α-type, and two concentrations of the crystalline polymorph of the metastable β-type due to strain-induced with a chain take planar zigzag structure having a chain exhibiting a 2 ₁ helical structure, affected by the collection method Can be done. These findings have a wide range of implications, as the crystalline structure of the polymer plays an important role in its properties developed after processing. In order to further elucidate the crystallization behavior of polymer chains during the electrospinning process, it is essential to study the internal structure of single electrospun nanofibers, including the spatial distribution of polymorphs. Unfortunately, there are very few techniques that can simultaneously provide the required spatial resolution and phase sensitivity / specificity.

The combination of atomic force microscopy (AFM) and infrared (IR) spectroscopy can overcome these technical limitations. This new technique, known as AFM-IR, is based on the photothermally induced resonance effect (PTIR). It is a powerful tool that provides topographical information that can correlate with chemical, structural and molecular orientation information with spatial resolution of 50-100 nm. Unlike traditional FT-IR spectroscopy, AFM-IR technology uses a sharp gold-coated AFM tip to detect the rapid thermal expansion of the sample caused by the absorption of short (10 ns) pulses of IR radiation. .. When the monochromatic laser radiation approaches the IR frequency that excites the molecular vibrations in the sample, the light is absorbed and induces rapid thermal expansion of the sample in contact with the AFM tip. This results in simultaneous deflection of the AFM tip, resulting in a "ring down" of the cantilever at its natural deflection resonance frequency as the heat dissipates. These movements of the cantilever are "detected" by a second laser beam reflected at the top of the cantilever, and this signal is measured using a position-sensitive photodetector. The resonant amplitude induced in the cantilever is proportional to the amount of IR radiation absorbed by the sample. Therefore, the final AFM-IR spectrum is obtained by measuring the ringdown amplitude while adjusting the IR laser on the IR fingerprint region. Further details regarding this AFM-IR device can be found elsewhere. The development of AFM-IR technology with the fine spatial resolution provided by the AFM tip and the phase sensitivity provided by IR spectroscopy is described herein to the spatial distribution of crystalline polymorphs in single electrospun nanofibers. Used for exploration.

In the present specification, a study of polymorphic distribution in single field-spun nanofibers using AFMIR technology is disclosed. Bio-based PHBHx nanofibers were produced by electrospinning on the tapered edges of a high speed rotating disk. The coexistence of α-type and β-type crystal polymorphs in single nanofibers was confirmed by both AFM-IR and selected area diffraction (SAED) by low-dose TEM. In addition, the dependence of β content and molecular orientation on fiber size was investigated by these two techniques on a single fiber scale. In addition, the spatial distribution of the two polymorphs over individual fibers of various diameters was examined by AFM-IR spectroscopy and imaging at 50 nm spatial resolution.

Further, in the present specification, a method for forming a β structure in a PHBHx film by stress-induced β crystallization is disclosed, and stress is applied by mechanically pulling the PHBHx film.

Further, the present specification discloses a method for forming a β structure in a PHBHx film by a novel method of mechanical tension following room temperature isothermal crystallization. It was confirmed that α crystallites must be present before the formation of the β structure. The ideal conditions for the initial formation of the β structure corresponded to 28 minutes of isothermal crystallization. Furthermore, it was shown that this β-structure is reversible and that the tensile process is different in the initial β-forming and re-tensile processes. The β structure can also be annealed to the α structure again at a temperature as low as 48 ° C.

Given the apparent importance of the β-type of PHBHx, it is desirable to ensure that this crystalline morphology is understood. Specifically, assessing the process by which the polymer transitions to this phase can facilitate the design of processing methods that promote β formation. IR and Raman spectroscopy are preferred analytical methods for this process given the number of vibration bands due to type β. In addition, the resulting spectrum can be easily analyzed using 2DCOS, which can determine the order of changes in the sample. The most promising means is to record the spectrum as a function of percent strain of PHBHx relative to the polymer membrane. Each spectroscope has its own limitations that need to be considered. In the infrared, it is necessary to measure the sample in transmission mode to prevent the membrane from relaxing. However, permeation measurements require a thin sample, which makes handling and tension difficult. Raman spectroscopy, on the other hand, does not require a thin sample, but the spectrum is recorded only at the laser spot and can be focused outside the strained region. To overcome this limitation, the membrane can be cut out to form a dogbone shape, which forces the sample to be distorted in that region. Therefore, Raman may have advantages beyond infrared. The resulting sample can be measured by XRD and DSC if possible. The XRD profile can be obtained with the sample remaining strained by positioning the mechanical tensioner within the instrument and thus keeping the sample tense. However, DSC analysis typically requires the sample to be removed from the mechanical tension device. A membrane that is stressed and annealed at a temperature low enough to increase the α content but not melt the β type may allow retention of the distorted sample. Preliminary results indicate that it may be preferable to insert the entire mechanical tensioner into a convection furnace and anneal the membrane for about 4 hours. Since the annealing temperature depends on the 3HHx content, a series of samples with β content can be annealed at different temperatures and then measured by XRD in order to properly determine the correct annealing temperature. Once the β peak disappears on XRD, a temperature below the annealing temperature of the sample may be selected to fix the strain structure. Although DSC measurements can be made with these samples, it should be noted that heating the samples can result in mechanical relaxation. To avoid this, the sample can be wrapped in a piece of aluminum prior to insertion into the DSC pan. Such analysis not only reveals a particular vibration band of type β and its formation process, but also identifies the thermal behavior of β crystals, i.e. the melting point. With this information, samples with higher β content can be designed more easily and applications can be created with thermal limitations in mind. Example 2 later disclosed herein provides an exemplary scrutiny of the polymorphic distribution in individual electrospun PHBHx fibers, including experimental procedures, results and considerations.

Devices Based on PHA-Based Copolymers In one aspect, the poly (hydroxybutyrate) PHA copolymer-based piezoelectric materials according to the invention are not limited to these, but for any suitable application including general purpose sensors, actuators, biobatteries, nanomotors. Can be used. When used in nanofiber form, any change in dimensions caused by, for example, humidity, temperature, salt content, attachment or injection of nutrients, attachment of metalloids, etc. will cause a potential difference (voltage) between its ends. Form. This produces a signal that can be detected remotely. There is no such multipurpose sensor, and the development of such a sensor is innovative and affects many environmental situations.

In another embodiment, the biocompatible, biodegradable and piezoelectric PHA copolymer-based piezoelectric materials according to the invention are used in disposable nanomotors and sensors. In one embodiment, the PHA copolymer-based piezoelectric material according to the present invention can be used for medical sensors and nanomotors.

FIG. 37 shows an exemplary device 100, wherein the device 100 comprises at least one ribbon or polarized polymer composition of electrospun fibers of the polyhydroxyalkanoate-based copolymer obtained by the method of claim 1. It comprises a PHA-based copolymer layer 110 comprising, the layer 110 being configured to exhibit one or more of a piezoelectric effect, a pyroelectric effect and a strong dielectric effect, a ribbon of electrospun fibers and a polarized polymer composition, respectively. Contains β-type PHA-based polymers present in an amount of about 10% to about 99% as measured by X-ray diffraction.

In one embodiment, the device 100 can be a sensor configured to generate a potential difference or voltage corresponding to the dashed line 120 in response to changes in the dimensions of the PHA-based copolymer layer. As shown in FIG. 37, in sensor applications, the change in dimensions of the sensor PHA-based copolymer 110 is the signal output (ie, charge or voltage) indicated by the broken line 120, where the arrow points in the opposite direction of the copolymer 110. To form.

In another embodiment, the device 100 may be an actuator configured to change dimensions (eg, expand or contract) in response to an application of charge or voltage corresponding to the dashed line 120 across the PHA-based copolymer layer. As shown in FIG. 37, in actuator applications, the charge input indicated by the dashed line 120, where the arrow points in the direction of the copolymer 110, causes a dimensional change in the PHA-based copolymer 110, which dimensional change is the desired effect. Is activated.

In the particular example schematically shown in FIG. 38, the sensor 100/230 may include a membrane / fiber 110 placed under tension. The presence of a stimulus (eg, adsorbed moisture in the fiber) can result in a change in membrane / fiber size 220 (eg, swelling), which creates a potential difference due to piezoelectric activity (210), which is digitally detected. Or converted to an analog signal (eg, lighting a light bulb) (240). Changes in other properties (eg, temperature, salt content, nutrient attachment or injection, metalloid attachment, etc.) can produce similar results and outputs.

In another particular example schematically shown in FIG. 39, the sensor 100/340 is in an implementation where the sensor is, for example, a "biobattery" (eg, an energy storage device driven by an organic compound, such as glucose detection). , Can be activated by biological stimuli. In the presence of the stimulus, the signal 120 read by a processor (eg, microcontroller 330) and communicated by a communication interface that can be wired or wireless (eg, by Bluetooth® technology) is a treatment device (eg, insulin pump 310). ), Which can then regulate the relevant parameter 350 (eg, increase or decrease insulin). In monitoring applications, the signal 120 emitted by the piezoelectric sensor 100 may simply be used to monitor the components of the patient environment (eg, in-situ glucose monitor). Implementation is not limited to the detection of glucose or the detection of components in blood. In another implementation, the biological stimulus may be a "heartbeat" produced by a physiological signal of an animal or human patient, such as a vasculature component (eg, heart, vein, artery), such as a heartbeat. Changes in the dimensions of the piezoelectric polymer due to (eg, local changes in pressure) are read by the processor 330 to generate a potential difference or voltage that is communicated (320) to the monitoring or treatment device 310. An exemplary treatment device 310 may include, for example, a pacemaker, which may adjust the pacemaker signal 350 depending on the measurement result. An exemplary monitoring device in this example may include, for example, a blood pressure monitor. Implementation of such sensors is not limited to any particular system of the animal or human body.

In another embodiment, the polymer products discussed herein (eg, electrospun fibers or other structures) are disclosed, for example, in US Pat. No. 9,897,547, which is incorporated herein by reference. As such, it can be configured for use as a sensor by the metallization of the structure and the addition of trapping molecules.

More specifically, the following represents specific embodiments of the present invention.

1. 1. A polarized material that is inherently stable up to the near crystal melting temperature range of the polar crystals of the material, or to the near softening temperature range or the glass transition temperature range of the material if the material is non-crystalline.

2. 2. The polarization material of Embodiment 1 in which an amount of plasticizing solvent that can be removed without substantial loss of polarization stability of the polarization material is distributed.

3. 3. The polarization material of Embodiment 1, wherein the dielectric constant is substantially increased by the presence of a solvent that improves the dielectric constant, as compared with the dielectric constant of the polarization material that does not contain the solvent.

4. The polarization material of Embodiment 3, which increases the dielectric constant by at least 50 percent.

5. The polarized material of Embodiment 1, wherein the material is a polarized poly (hydroxybutyrate-co-hydroxyhexanoate) (PHBHx) copolymer, wherein the polarization is essentially stable up to the crystal melting range of the polar crystals.

6. The polarized material of Embodiment 1, wherein the material is a polarized PHBHx copolymer.

7. The polarized material of Embodiment 1, wherein the material is a polarized poly (hydroxybutyrate) / poly (vinylidene fluoride) copolymer.

8. One selected from the group consisting of soluble ceramic materials, PHBHx copolymers, poly (vinylidene fluoride), vinylidene copolymers, nylon-3, nylon-5, nylon-7, nylon-11, and blends thereof. The polarization material of Embodiment 1, which comprises a plurality of components.

9. The material is a polymer having polymer units that can be polarized, hydroxybutyrate units, hydroxyhexanoate units, vinyl units, vinylidene units, ethylene units, acrylate units, methacrylate units, nylon units, carbonate units, acrylonitrile units, cellulose. 1 selected from the group consisting of esters other than esters of units, pendant fluoros, chloros, amides, acrylates and methacrylates, cyanides, nitriles other than acrylonitrile units, or units having ether groups, protein units, or combinations thereof. The polarization material of Embodiment 1, comprising one or more components.

10. A method for preparing a polarization material,
a) To provide a melt composition of one or more polymers in the form of a membrane that can be polarized in a preselected form;
b) With the step of quenching the molten polymer composition;
c) With the step of cold stretching the polymer composition;
d) A step of polarizing a cold-stretched polymer composition by applying an effectively high electric field with a strength lower than the strength that results in substantial dielectric breakdown of the polymer material;
e) A step of annealing a polarized polymer composition at an annealing temperature lower than the melting temperature of the crystals of the polarized polymer composition, whereby polarization is maintained and polarization is performed at or near the crystal melting point of the polar crystals of the polymer composition. A method comprising a step in which thermal stability is provided.

11. The method of embodiment 10, wherein the composition comprises a solution of one or more polymers and one or more solvents for the polymers.

12. The polymer has a polymer unit that can be polarized, and has a hydroxybutyrate unit, a hydroxyhexanoate unit, a vinyl unit, a vinylidene unit, an ethylene unit, an acrylate unit, a methacrylate unit, a nylon unit, a carbonate unit, an acrylonitrile unit, and a cellulose unit. , Pendant fluoro, chloro, amide, acrylate and ester other than the methacrylate unit ester, cyanide, nitrile other than the acrylonitrile unit, or a unit having an ether group, a protein unit, or a combination thereof. Alternatively, the method of embodiment 10 comprising a plurality of components.

13. The composition comprises polyvinyl chloride, polymethylacrylate, polymethylmethacrylate, poly (vinylidene cyanide / vinyl acetate) copolymer, vinylidene cyanide / vinyl benzoate copolymer, vinylidene cyanide / isobutylene copolymer, vinylidene cyanide / methyl methacrylate. Implementations comprising one or more components selected from the group consisting of copolymers, polyvinyl fluoride, polyacrylonitrile, polycarbonate, cellulose, proteins, synthetic polyester and ethers of cellulose, and poly (γ-methyl-L-glutamate). Form 10 method.

14. The method of embodiment 10, wherein the composition comprises one or more components selected from the group consisting of poly (vinylidene fluoride) and vinylidene fluoride copolymers.

15. The composition is poly (hydroxybutyrate / hydroxyhexanoate), poly (vinylidene fluoride / trifluoroethylene), poly (vinylidene fluoride / tetrafluoroethylene), poly (vinylidene fluoride / vinyl trifluoride). , The method of embodiment 14, comprising one or more copolymers selected from the group consisting of poly (vinylidene fluoride / vinyl chloride) and poly (vinylidene fluoride / methylmethacrylate).

16. The composition is selected from the group consisting of soluble ceramic materials, PHBHx, poly (vinylidene fluoride), vinylidene copolymers, nylon-3, nylon-5, nylon-7, nylon-9 and nylon-11 and blends thereof. The method of embodiment 10 comprising one or more components.

17. The method of embodiment 10, wherein the composition comprises nylon-7 and nylon-11.

18. The method of embodiment 10, wherein the composition comprises nylon-7 and nylon-11 by weight of about 50:50, respectively.

19. The method of embodiment 10, wherein the composition comprises a melt or polymer of the polymer.

20. The method of embodiment 10, wherein the membrane comprises at least two layers of one or more polymers that can be polarized, formed from the polymorphic composition of the polymer.

21. 10. The method of embodiment 10, wherein the membrane comprises at least two layers of one or more polymers that may be co-extruded and in contact with each other and polarized.

22. The method of embodiment 10, wherein the membrane comprises a self-supporting sheet of the composition.

23. The method of embodiment 10, wherein the film comprises a non-self-supporting layer of the composition placed on a self-supporting sheet.

24. The method of embodiment 10, wherein the polarization is carried out at about 20 ° C. to about 120 ° C. for up to about 5 hours using an electric field of at least 1 MV / cm.

25. The method of embodiment 10, wherein the membrane is annealed at a temperature of about 125 ° C to about 150 ° C for at least 1 hour.

26. The method of claim 10, wherein the composition comprises a blend of PHBHx, poly (vinylidene fluoride) and vinylidene fluoride-vinyl fluoride (80/20) copolymer.

27. The method of embodiment 10, wherein the composition comprises a blend of PHBHx at 50:50 by weight and a vinylidene fluoride-vinyl fluoride (80/20) copolymer.

28. The method of embodiment 10, wherein the composition comprises a blend of PHBHx and poly (vinylidene fluoride) at 50:50 by weight.

Example

Forming β-type crystal structure of electrospun poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx) nanofibers Material 3.9 mol% Hx content Poly (3-hydroxybutyrate-co-3-hydroxyhexanoate) (PHBHx) (Mw = 843000 g / mol, PDI = 2.2) having the above was supplied from Crystal & Gamble Company. The polymer was purified by dissolving it in chloroform, followed by filtration and then precipitating in hexanes. Solvents 1,1,1,3,3,3-hexafluoro-2-propanol (HFIP) were purchased from Sigma-Aldrich and used as supplied.

Electrospinning The purified PHBHx was dissolved in HFIP and stirred at 60 ° C. overnight to ensure complete dissolution to prepare a 1 wt% electrospinning solution. As part of an experimental protocol for electrospinning nanofibers, a 3 mL BD plastic syringe equipped with a 21 gauge stainless steel needle was filled with a polymer solution and connected to a high voltage power supply positive electrode held at 10 kV. .. Two different negatively charged collectors, a parallel electrode collector and a rotating disc collector were used to collect electrospun nanofibers with the desired morphology. In the parallel electrode collector, a rectangular slot was cut into the aluminum foil piece, and the slot was made into an air gap of 35 mm × 10 mm. In a rotating disc collector, the disc was designed to have a tapered edge with a 30 ° half-width to form a convergent electric field. The angular velocity of the rotating disk was set to 3500 rpm, which corresponds to the linear velocity of 1117 m / min at the edge of the disk. The applied voltage between the needle and the collector was 25 kV. Working distance and solution feeding rate were 25 cm and 0.5 mL / hour, respectively. All electrospun mats were dried in vacuum for 24 hours to remove any residual solvent prior to further investigation.

Characteristic determination Fiber mat: Using a field emission scanning electron microscope (SEM, JEOL JSM 7400F), the morphology of field-spun PHBHx nanofibers was observed at an acceleration voltage of 3.0 kV. Fiber diameters were measured using ImageJ software. Wide-angle X-ray diffraction (WAXD) measurements were performed under ambient conditions using a Cu Kα (λ = 1.5418 Å) as the X-ray source and a Rigaku Ultima (IV) instrument operating at 44 kV and 40 mA. Scanning was performed at a rate of 1 ° / min and a sample step of 0.1 ° within a 2θ range of 10 ° to 40 °. A Fourier Transform Infrared (FTIR) spectrum was collected using Thermo Nicolet NEXUS 670 in transmission mode at room temperature. For each sample, 128 scans were signal averaged with a spectral resolution of ^{4 cm -1.}

Single fiber: Selected area diffraction (SAED) patterns and brightfield images were recorded by a transmission electron microscope (TEM, Technai G2 12) equipped with a low dose CCD camera at an acceleration voltage of 120 kV. Nanofibers were deposited on a 300 mesh copper grid coated with a lace-like carbon film to reduce specimen damage. When performing the SAED experiment, the diffraction pattern was acquired with a fixed camera length of 2.1 m, and the TEM image was taken at a constant magnification of 97,000. Prior to fiber deposition, a thin layer of gold polycrystal was sputtered onto each of the copper grids, which was used to calibrate camera constants and correct distortions in any system.

High resolution AFM images and IR spectra of single electrospun fibers were acquired by NanoIR2 AFM-IR (Anasis Instruments). PHBHx nanofibers were electrospun directly onto a transparent silicon wafer substrate in the mid-infrared region of ^{900-3600 cm-1} to maintain good contact between the sample and the substrate. NanoIR spectra were collected with a spectral resolution of 2 cm ^-1 and simultaneous averaging of 256 cantilever ringdowns for each data point.

Results and Discussion Studies on Fiber Mats: In recent years, many studies have described the significant effects of collectors on fiber morphology, macroscopic alignment, and molecular orientation. The morphology of electrospun PHBHx nanofibers collected using different collectors while maintaining the same other electrospinning parameters was examined by SEM, the image of which is shown in FIG. The SEM image clearly shows how the collector affects the fiber morphology. Due to bending instability in the whipping region, the fibers collected on the aluminum foil outside the gap are random (Al foil random fibers, FIG. 1a), while the fibers collected over the air gap (air gap aligned fibers, The fibers collected on the tapered edges of the rotating disc (FIG. 1b) and the rotating disc (rotating disc aligned fibers, FIG. 1c) are well aligned. In addition, the average diameter of the aligned fibers (270 ± 20 nm) was much smaller than the average diameter of the Al foil random fibers (500 ± 30 nm), indicating additional tension and stretching during formation. .. For a long time, in rotary collection, it was recognized that the electrospun fibers were further pulled and aligned in the winding direction. However, in the case of air gap aligned fibers, the tension is provided by the electrostatic attraction between the residual positive charge on the fiber and the negative charge accumulated on the gap edge. These gravitational forces, in coordination with the repulsive force between the residual charge on the unreleased fibers in the air gap, result in a macroscopic alignment of the fibers.

In this study, the molecular orientation in the fiber mat was characterized by polarized Fourier transform infrared spectroscopy (p-FTIR). P-FTIR has been widely used to study molecular orientation and structural changes in polymer chains during electrospinning. In an FTIR spectrum polarized in a particular direction of an incident infrared beam, high absorbance is measured if the change in the dipole moment of vibration has a component along the electrical vector of the incident beam. FIG. 2 shows parallel and vertical p-FTIR spectra of Al foil random fibers (FIG. 2a), air gap aligned fibers (FIG. 2b), and rotating disc aligned fibers (FIG. 2c). As shown in the figure, the p-FTIR spectra of the two aligned fibers in perpendicular directions showed a clear difference in absorbance, indicating the presence of molecular orientation of the PHBHx polymer chain. To quantify this microscopic orientation, the following equation (assuming uniaxial symmetry) was used to calculate the normalized dichroism differences (NDDs) for some characteristic peaks.

In the equation, A‖ is the parallel polarized infrared absorbance with respect to the macroscopic fiber axis, and A⊥ is the vertically polarized infrared absorbance. As can be seen from this equation, NDD is in the range of -1 / 2 to 1, and when the sample is isotropic, NDD = 0. As listed in Table 1, the NDDs for carbonyl expansion and contraction were calculated to be 0.001, -0.134 and -0.100 for Al foil random fibers, air gap aligned fibers and rotating disc aligned fibers, respectively. .. When NDD <0, the absorbance of C = O is lower when the electrical vector of the incident infrared beam is parallel to the fiber axis than when the vector is perpendicular to the fiber axis (A‖). <A⊥) Indicates that. Since the carbonyl bond is approximately perpendicular to the molecular skeleton as predicted by its chemical structure, this result shows that the carbonyl bond is oriented perpendicular to the fiber axis and the polymer chain is along the fiber axis. It was suggested that it was oriented. As the NDD of the carbonyl bond approaches -1 / 2, the chains are more oriented along the fiber axis. On the other hand, NDD> 0 of the C—O—C bond along the molecular skeleton suggests that the C—O—C bond was oriented almost parallel to the fiber axis. As the NDD of the C—O—C bond approaches 1 (1.0), the chains are more oriented along the fiber axis. It should be noted from Table 1 that the airgap aligned fibers had the lowest NDD (highest absolute value) for C = O stretch and the highest NDD for C—O—C stretch. This suggests that the air gap aligned fibers had the highest level of chain orientation along the fiber axis. In particular, it should be noted that the NDD of the three bands was always around 0 in the case of Al foil random fibers. This result is due to the lack of macroscopic alignment of the fibers. Under these conditions, nothing can be said about the degree of chain alignment within the individual fibers.

The crystal morphology of electrospun PHBHx nanofibers was examined by WAXD and FTIR. FIG. 3 shows the WAXD profiles of Al foil random fibers, air gap aligned fibers, and rotating disc aligned fibers. In all profiles, diffraction peaks attributable to α-type having a 2 ₁ helical structure was observed. Furthermore, in the WAXD profile of the air gap aligned fiber and the rotating disk aligned fiber, a diffraction peak attributed to β type having a planar zigzag chain structure was observed at 2θ = 19.6 °. As is clear, the rotating disc alignment fibers had significantly more β-type than the air gap alignment fibers. For display purposes, the intensity of α (020) at 2θ = 13.7 ° in all profiles was normalized to 1. Considering that β-type was not found in Al foil random fibers, different collection strategies (air gap vs. rotating disc rotating at 3500 rpm) were generated with substantially different tensile forces on the fibers, and the corresponding latter. Due to the increased amount of β-type produced, this elongated chain structure should be induced by two improved collection methods that yielded macroscopically aligned fibers.

[Table 1]
Table 1: Normalized dichroism differences (NDDs) for different vibration bands

The introduction of a planar zigzag chain structure in the fibers collected over the gap or on the rotating disk was further confirmed in the transmissive FT-IR spectrum (FIG. 4). C = O expansion and contraction in PHA strongly correlates with the polymer backbone structure. As can be seen in FIG. 4a, the C = O stretch band of the out-gap random fiber can be decomposed into ^{a strong peak of 1725 cm -1 corresponding to the crystalline phase and a weak shoulder of 1746 cm -1} corresponding to the amorphous phase. As the amount of β-type increases from random fibers to aligned fibers (normalized WAXD results), ^{the peak of 1725 cm -1} becomes wider and shifts to higher frequencies, while ^{the shoulder of 1746 cm -1} has its relative strength. Increased and shifted to lower frequencies. In other words, as the concentration of β crystal morphology increased, the peaks and shoulders gradually approached each other and became more similar in shape. In addition, some spectral changes also occurred in the fingerprint region. ^{In FIG. 4b, the peaks of 1304, 1080, and 969 cm -1} increased with the increase of β-type, while the peaks of 1277, 1047, and 950 cm ^-1 decreased and became shoulders. Similar findings have been obtained in other studies. These findings are important because they show a correlation between the presence of metastable β-structures (confirmed by WAXD) and changes in the vibrational spectrum, especially the appearance of new bands. As a result, the above peaks can be considered to indicate the presence of planar zigzag skeletal properties of β-crystal polymorphs.

Studies on single fibers: Selected area electron diffraction (SAED) with low dose TEM was used to examine crystal structure and orientation on a single fiber scale. FIG. 5 shows individual electrospun PHBHx fibers collected on aluminum foil (a, a'), over the air gap (b, b'), and on the tapered edges of the rotating disk (c, c'). A bright field image and a SAED pattern are shown. The fibers from the air gap (b) and the rotating disk (c) corresponded to diameters of 195 nm and 221 nm, respectively, while the fibers from the aluminum foil (a) were twice as large and had a diameter of 495 nm. The corresponding SAED patterns of these three fibers showed distinct differences from each other, and all three SAED patterns had crystal reflections characteristic of orthorhombic α-type crystals, but crystals when different collectors were used. The orientation of was significantly changed. In this study, the crystal orientation is tangentially widened in electron diffraction, or equatorial α (020) and α (110), meridional α (002), and layer α (111), as summarized in Table 2. It was quantified by the center angle (ψ) of the arcuate reflection. The central angle of the arc corresponds to the angular distribution of the indexed crystal plane, so a smaller ψ indicates a higher degree of uniaxial orientation of the crystal. It is noted from Table 2 that the fibers from the rotating disc always showed the smallest ψ for all arcs, which is the highest of the three samples in this single fiber. It has a degree of orientation, suggesting that the polymer chains in the crystal were highly oriented along the fiber axis. Crystal reflections of β-type crystals appear in the fibers from the two improved collectors, and new equatorial arc pairs are observed in both 4b'and 4c', which are attributed to the β-type crystal plane. Since the β-type crystal structure is a strain-induced quasicrystal structure, this finding does indicate that the two improved collectors provide an additional tensile force sufficient to pull the polymer chain to an essentially fully extended form. It indicates that it was introduced into the fiber. It is also noted that the β-type arc always had a small central angle of 8 ° (Table 2), which is the molecule in the β crystal along the fiber axis under different collection conditions. It suggests a high degree of chain orientation.

[Table 2]
Table 2: Outline of SAED pattern of single fiber collected by different methods and central angle of arc reflection of indexed crystal face.

For each of the three SAED patterns, the intensity profile along the equator line is plotted against the scattering vector (1 / d, the reciprocal of the spatial distance), and the obtained profile is shown in Table 2. In order to eliminate the influence of non-structure related factors such as electron beam conditions and optical recording conditions, the gray values of all backgrounds of the three SAED patterns were equalized prior to the analysis. The baseline was corrected by fitting the background to each of the profiles. The following findings were obtained from these intensity profiles: First, from left to right, the intensity of both the α (020) and α (110) peaks decreased while the intensity of the β peak increased, which correlates with the decrease in the α crystal structure content. It shows an increase in β crystal structure content. These findings are due to the electrostatic attraction between the positively charged fibers and the negatively charged gap edges, and the tensile force from the insulating gap caused by the electrostatic repulsion between the residual positive charges on each fiber. It shows that it is significantly weaker than the tensile force from the rotating disk. Furthermore, the correlated increase in β crystal content with the decrease in α crystal content is such that the formation of β-type crystal structure is initiated by the tensile force from the two improved collectors, and the α-type and β-type crystal structures are It suggests that it was formed simultaneously from amorphous and mobile polymer chains by two competing crystallization processes during collection. It should be noted that the formation mechanism of this β-type crystal structure is different from that proposed by Iwata and Ishii, in which β crystals are formed after α crystals. Second, the full width at half maximum (fhm) of the two α peaks increased, indicating a decrease in effective crystal size according to Scherrer's equation. However, the fwhm of the β peaks in the last two profiles was similar, suggesting that the effective size of the β crystals remains the same under different tensile forces. The decrease in the effective size of the α crystal may be due to the faster solvent evaporation and thus the faster solidification when using the improved collector. As a result, the polymer chains were fixed in the initial state of crystallization. That is, the additional tensile force provided by the improved collector during the collection process initiates the formation of the β crystal structure by extending the mobile amorphous chain into a planar zigzag structure, which is the formation of the α crystal structure. Competing. In addition, these tensile forces improve the orientation of the α crystal along the fiber axis and reduce its effective size. However, the effect of tensile force on the degree of orientation and size of β crystals is limited.

Careful examination of the SAED patterns of 5b'and 5c' shows that, especially in 5b', two arcs with different widths in the azimuth direction for each indexed α crystal plane listed in Table 2. It was found that there was always an overlap of (see the outline in Table 2). One arc had a larger central angle but was narrower, and the other had a smaller central angle but was wider, which had two sets of α crystals with different degrees of orientation and crystal size. It is shown that. It was also observed that as the tensile force increased from 5a'to 5c', the narrower arcs became much smaller, while the wider arcs became only slightly smaller, with two sets of tensions. It is shown that it had a different effect on the orientation of α crystals. Another interesting observation was the appearance of a large α (001) layer line over the meridian at 5b', while two pairs of different α (011) arc reflections were observed at 5c'. These results show the change in the filling state of the molecular chain in the α-type crystal structure along the fiber axis.

In summary, the results from SAED experiments on single fibers support the significant effect of collection methods on crystal structure and crystal orientation levels, which is consistent with the conclusions drawn from studies on fiber bundles. In addition, SAED results also demonstrated that substantially polymer chain orientation also occurred in randomly collected fibers on aluminum foil, which could not be determined in previous studies on fiber bundles. .. These results further indicate that the tensile force during the electrospinning process is large enough to partially orient the chain, but not large enough to extend the chain into a planar zigzag structure, which is additional during collection. It further shows that tensile force is required. In addition, in SAED experiments, rotating disc aligned fibers were observed to have the highest level of chain orientation. This is in contrast to the results obtained from polarized FT-IR experiments in which the bundles of airgap aligned fibers clearly show the highest levels of chain orientation in all three samples. This difference may be due to inconsistencies and size / morphological inhomogeneities (beads) of rotating disc aligned fibers within the bundle resulting in averaging of polarized FT-IR signals across the fiber bundle. As a result, the investigation of individual polymer nanofibers becomes increasingly important.

As mentioned above, SAED is a powerful technique for investigating single electrospun nanofibers. However, SAED experiments are sometimes difficult and time consuming. To complement these results, we used AFM-IR, a novel technique that allows direct investigation of both crystalline and amorphous phases of ultrafine electrospun fibers on a single fiber scale. AFM-IR is a technique that combines atomic force microscopy (AFM) and infrared spectroscopy (IR) for nanoscale characterization. It simultaneously provides IR spectra and AFM images with features less than 100 nm. The light source is an adjustable IR laser whose wavelength can be swept over the infrared "fingerprint" region in less than a minute. When one of the wavelengths is absorbed by the sample, thermal expansion of the sample occurs on a nanosecond time scale, which results in modulation of the oscillating AFM cantilever. This causes a "ring down" at that particular frequency, which diminishes as the heat dissipates. The positive amplitude of the vibration represents the IR band intensity, so when the frequency is ^{adjusted over the IR region (900-3600 cm-1} ), the IR spectrum is obtained with spatial resolution of 50-100 nm. Further details of this device have been reported elsewhere.

To test the feasibility of this technique, IR spectra of a single Al foil fiber and a single rotating disc fiber were collected and compared to the transmissive FT-IR spectra of the corresponding fiber mats (FIG. 3, FIG. Black and blue spectra). 6a and 6b are AFM images of two fibers having diameters of 473 nm and 324 nm, respectively. The IR spectra of these two single fibers are shown in FIG. 6c. As shown in the figure, the spectrum of fibers collected on the Al foil decomposes into two features, a strong peak of 1725 cm ^{-1 and a weak shoulder of 1746 cm -1, corresponding to the α crystalline phase and the amorphous phase, respectively.} Could be done. However, in the spectrum of fibers collected on the rotating disk, ^{two different peaks of 1728 cm- 1} and 1740 cm-1 were observed, which were assigned to the more regular α and β crystal phases, respectively. These nano-IR spectra are in good agreement with conventional FT-IR spectra in terms of peak / shoulder position and relative intensity. However, it should be noted that the peaks in the nano-IR spectrum are more highly resolved compared to the corresponding FT-IR spectra, which are broadly blurred due to averaging over the distribution of slightly unmatched fibers. ..

Mechanism of formation of β-type crystal structure The formation of β-type crystal structure can have a significant effect on various properties of the material. To date, this strain-induced metastable crystal structure has been characterized by different modes of processed PHBs and PHBVs, including hot / cold stretched films after isothermal crystallization, two-step stretched fibers, and one-step stretched fibers. It has been reported in crystallization materials. In these highly pulled PHB or PHBV thin films and fibers, β-type was thought to result from free chains in the amorphous phase between well-developed α-layered crystals. In other words, the β-type crystal structure is generated after the formation of α crystals. However, due to the large amount of amorphous chains in the material that cannot be highly elongated during processing, PHBHx52 cannot obtain a β-type crystal structure by using the same processing method.

In the present disclosure, the β-type crystal structure in PHBHx was successfully generated by collecting nanofibers on a high-speed rotating disk, but the crystallinity of the obtained fibers is suggested by preliminary DSC measurement. It is as low as 44 ± 1% (when the crystallinity is calculated by Xc = ΔHm / ΔHm0 × 100%, where ΔHm0 is the melt enthalpy (146J / g53) of the 100% crystalline PHB homopolymer. ). Experimental that α-type and β-type crystals coexist in the resulting fiber at the same time, and an increase in β crystal content correlates with a decrease in α crystal content due to an increase in tensile force during collection (see SAED). Based on observations, it can be concluded that both α and β crystal morphologies are formed during collection. Both appear to be formed from amorphous and mobile chains by two different competing crystallization processes. The possible formation mechanism of the β crystal structure is shown in FIG. Unlike the previously reported formation mechanism (β crystals are formed after α crystals), the β type in electrospun PHBHx nanofibers is actually at the same time as the formation of α type (path 1) or even itss. It is formed before (path 2). The melted polymer chains in the random coil state are highly pulled during electrospinning and are therefore elongated and oriented along the tensile direction. As a result, the amorphous fibers that reach the air gap or rotating disc consist of polymer chains oriented along the fiber axis and are plasticized by the residual solvent. During collection, the amorphous fibers are further pulled by pulling over the gap or by the additional stretching force provided by hoisting to a high speed rotating wheel. At the highest elongation, some oriented chains in the amorphous fiber are fully elongated and have a planar zigzag structure. At high solvent evaporation rates, this metastable β structure is immobilized within the solidified fibers, so β crystals are formed at the same time as the α crystals, as shown in Path 1. As a further possibility, under strong tension during collection, most of the polymer chains in the amorphous fiber are fully elongated and have a planar zigzag structure. At lower evaporation rates due to the capture of residual solvent in the core of the fiber immediately after spinning, some of the chains, especially within the core, are relaxed and converted to a more stable helical structure (alpha type). Therefore, the final solidified fiber contains both α and β crystal structures, as shown in Path 2. In reality, the actual formation mechanism can be a combination of both, and hopefully ongoing studies will provide further insight into the process of β-form formation in PHBHx.

The invention is exemplified and described herein with reference to specific embodiments, but the invention is not intended to be limited to the details presented. Rather, various modifications may be made in detail within the scope of the claims and their equivalents without departing from the present invention.

Conclusion We investigated the fine structure of electrospun PHBHx nanofibers for fiber mats and single nanofibers. The molecular chain structure, crystal structure, and crystal / chain orientation were investigated by polarized FT-IR, WAXD, SAED, and AFM-IR and found to be highly dependent on the collection method. More importantly, by using two improved collectors, strain-induced (strain-induced) β-type crystal structures were obtained for the first time in electrospun PHBHx nanofibers. The β type was identified based on the appearance of new crystal reflections in WAXD and SAED, as well as the spectral changes observed in the IR spectra. In addition, the results from SAED experiments on individual fibers provided insights into the correlation between tensile force and the degree of orientation and size of α and β crystals. Finally, AFM-IR technology has demonstrated to be a powerful and efficient tool for microstructural investigation of individual electrospun nanofibers. Furthermore, experimental results suggest a new mechanism for the formation of β crystal structures that is significantly different from previously reported, although we do not want to be bound by any particular theory. The β crystals resulting from the oriented free chains in the fibers were formed at the same time as or even before the formation of the α crystals during collection. This study has led to further investigation of the relationship between β-structure and the mechanical properties and processing protocols of PHBHx. The excellent biodegradability and biocompatibility of the material, as well as the corresponding changes in the macroscopic performance of the material, make PHBHx a promising material in many application areas.

Polymorph distribution in individual electrospun poly [(R) -3 hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx) Polymer 3.9 mol% hydroxyhexanoate (Hx) Poly (3-hydroxybutyrate-co-3 hydroxyhexanoate) (PHBHx) (Mw = 843000 g / mol, PDI = 2.2) produced by a bacterium having a comonomer content is supplied from Polymer & Gamble Company. Was done. The polymer was purified by dissolving it in chloroform (Fisher Scientific), followed by filtration and then precipitating in hexane (Fisher Scientific). Solvents for electrospinning 1,1,1,3,3,3-hexafluoro-2-propanol (HFIP) were purchased from Sigma-Aldrich and used as supplied.

Electrospinning The purified PHBHx was dissolved in HFIP and stirred at 60 ° C. overnight to ensure complete dissolution to prepare a 1 wt% electrospinning solution. At room temperature, a 3 mL BD plastic syringe equipped with a 21 gauge stainless steel needle was filled with a polymer solution and connected to the positive electrode of a high voltage power source held at 10 kV. Macroscopically aligned electrospun nanofibers were collected using a 5 mm thick negatively charged (-5 kV) rotating disc. The angular velocity of the rotating disk was set to 3500 rpm, which corresponds to the linear velocity of 1117 m / min at the flat edge. Working distance and solution pumping speed were 25 cm and 0.5 mL / hour, respectively. A transparent silicon wafer (Addison Engineering, Inc.) in the mid-infrared region (900 to 3600 cm ^-1 ) was cut into pieces of 5 mm (width) x 8 mm (length) and attached to the edge of a rotating disk. During electrospinning, the fibers were electrospun directly onto a silicon wafer to maintain good contact between the fibers and the substrate. The fiber density on the silicon wafer can be easily adjusted by controlling the length of the electrospinning time, and the electrospinning time is an approximate fiber density of 2 fibers / mm (length) in this study. Set to 45 seconds to get. After fiber deposition, the silicon wafer was placed in vacuum for 24 hours to remove any residual solvent prior to further investigation.

Cutting with a microtome A bundle of rotating disk aligned nanofibers was embedded in parallel in 2 Ton epoxy (ITW Devcon) prior to cutting with a microtome. After curing, 250 nm thick sections were cut out by a microtome (Leica Ultracut UCT) at room temperature. The thin sections were then transferred to 10 mm × 10 mm flat ZnS for the AFMIR test.

Selected Area Electrodiffraction (SAED)
SAED patterns and brightfield images were recorded by a transmission electron microscope (TEM, Technai G2 12) equipped with a low-dose CCD camera using an acceleration voltage of 120 kV. Nanofibers were deposited on a carbon-coated 400 mesh copper grid to reduce sample damage. Diffraction patterns were acquired with a fixed camera length of 2.1 m. Prior to fiber deposition, a thin layer of gold polycrystal was sputtered onto each of the copper grids, which was used to calibrate camera constants and correct distortions in any system.

AFM-IR measurement: spectroscopy vs. imaging. Nanoscale infrared measurements were performed using a nanoIR2 platform (Anasis Instruments, Santa Barbara, CA) that focused radiation from an adjustable IR laser source from above onto a location on the top surface of the sample. Fibers were inspected in contact mode using gold-coated SiN AFM tips with a nominal tip radius of 20 nm. In this study, the output of the incident IR laser was adjusted to about 2% of the open beam intensity in order to obtain a good ring-down signal. In addition, an additional mesh filter was placed in front of the IR laser to further attenuate the beam in order to avoid melting / softening of the sample. The IR mapping was adjusted to match the IR signal updates and the pixel rates of the images. AFM height and IR peak images were primary flattened using the equipment's embedded software (Analysis Studio, Analytics Instruments). AFM-IR spectra were collected at data point spacing of ^{2 cm -1,} were co-averaging the cantilever ring-down 256 in the spectral range 1680~1780cm ^-1. The actual spectral resolution within this wavenumber range is 4 cm ^-1 , which is the laser line width. For each sampling point on the fiber, five spectra from the same position were averaged to reach a satisfactory signal-to-noise ratio. All measurements were made under ambient conditions.

Results and Discussion By using a rotating disc as the fiber collector, macroscopically aligned electrospun PHBHx nanofibers could be obtained. In this study, individual PHBHx nanofibers from the same batch were tested by AFM-IR. 8a and 8a'show AFM height images of two single rotating disc aligned fibers. The two fibers differ in size and have diameters of 263 nm and 390 nm, respectively. Example 1 showed that when the fibers were collected on a high speed rotating disk, a strain-induced semi-stable β-type crystal structure in which the elongated polymer chains had a planar zigzag structure was introduced. This result suggests that under this particular electrospinning condition, the appearance of the β crystal phase is indicated by the characteristic IR absorption at ^{1740 cm -1 of the IR spectrum.} The IR absorption peak of 1728 cm- ¹ is characteristic of the α crystal phase, which is a thermodynamically stable crystal structure of PHBHx. To investigate the spatial distribution of the two crystalline phases of nanofibers, adjusts the frequency of the incident IR laser to 1740 cm ^-1 and 1728 cm ^-1 are two characteristic frequencies corresponding to the β-type and α-type crystal structure, respectively Thereby, two single fibers were imaged. IR images are shown in FIGS. 8b and 8b'. In each IR image, the upper half is recorded in 1740 cm ^-1, the lower half was recorded at 1728 cm ^-1. Careful examination of these IR images can yield some findings. First, a comparison of FIGS. 8b and Fig. 8b ', the fibers were observed to have a detectable IR absorption at both 1740 cm ^-1 and 1728 cm ^-1, which, alpha crystal phase in a single fiber scales and β It shows the coexistence of crystalline phases. Furthermore, the same mapping conditions, finer fibers the higher absorption (263 nm) is in 1740 cm ^-1, whereas a lower absorption at 1728 cm ^-1 in which, finer fibers are thicker fibers (390 nm) It shows that it tends to have a higher β content than. Since the two fibers were manufactured under the same electrospinning conditions, the difference in diameter should be due to the different tensile forces they received during electrospinning and collection. Assuming elongation strain, the draw ratio of the 263 nm fiber is about 2.2 times larger than the draw ratio of the 390 nm fiber. Since higher draw ratios favor elongated chains, more molecular chains in the finer fibers are found in a planar zigzag structure and are filled to form β crystal morphology.

The coexistence of both α-type and β-type crystal structures in a single PHBHx nanofiber, as well as the dependence of the β-type content on the fiber diameter, is reconfirmed by the SAED pattern of the individual PHBHx nanofibers. 9a and 9a'show the original SAED pattern of two single PHBHx nanofibers from the same batch. The inset shows a brightfield TEM image of two single fibers. The diameters of the two fibers are 251 nm and 417 nm, respectively, which are equivalent to the diameters of the two fibers in the AFM-IR study. For comparison purposes, the original SAED pattern was contrast inverted as shown in FIGS. 9b and 9b'. By comparing the two inverted SAED patterns, it is observed that the equatorial arcs assigned to the β-type crystal plane 14 are clearly shown (indicated by the red arrows) in FIG. 9b, while in FIG. 2b', they are. Almost unrecognizable. For each of the SAED patterns, the intensity profile along the equatorial line was plotted against the scatter vector (1 / d, reciprocal of spatial distance) in correspondence with FIGS. 9c and 9c'. Clearly, the intensity of the β peak in FIG. 9c is much higher than the β peak in FIG. 9c', indicating that the finer fibers contain more β-type crystals than the thicker fibers. It was also observed that the arc in FIG. 9b has a less tangential spread or a smaller central angle than the arc in FIG. 9b', which means that thinner fibers (251 nm) are thicker (417 nm). ), It shows that it has a higher molecular orientation. That is, both the content of the β-type crystal structure and the degree of molecular orientation in the single electrospun PHBHx nanofiber increase with the decrease in the fiber diameter. Based on the results of AFM-IR and SAED studies, a strong correlation between β-type content and fiber diameter in single PHBHx fibers was observed, which is the tension received during electrospinning and nanofiber collection processes. It depends on the force. The finer the fiber, the higher the β-type content and the higher the degree of molecular orientation.

The second finding that can be obtained from the IR images in FIGS. 8b and 8b'is that for both fibers ^{, absorption at 1740 cm -1} is always high along the fiber edge, especially in FIG. 8b'. It is almost completely concentrated on the two edges. This finding suggests an inhomogeneous spatial distribution of α and β crystal structures throughout the fiber. More importantly, this suggests an interesting core-shell structure containing β-crystal polymorphs with far more shells than cores. In addition, careful examination of the two IR images revealed that the shell thickness, indicated by the width of the red line along the fiber edges, was about 10 nm, independent of the fiber size.

To test the hypothesis of fiber core-shell structure, AFM-IR spectra were collected at different positions for each of the two fibers, but IR images show the presence of heterogeneity. The AFM-IR spectra are shown in FIGS. 8c and 8c'. All AFM-IR spectrum has two peaks characteristic of 1740 cm ^-1 and 1728 cm ^-1, which are respectively known to correlate to the carbonyl stretching of β crystalline form and α crystalline form (14 ). However, the relative strength of the two peaks depends on the fiber size and position within the fiber. For example, in FIG. 8c, the three spectra collected on the fiber edges E1, E2, and E3 have a relatively high 1740 cm- ¹ peak, while collected at the fiber centers C1, C2, and C3. The three spectra have a relatively high 1728 cm- ¹ peak. Similar findings were obtained in FIG. 8c'. For each fiber, spectra collected at different spots at the edges or centers have the same spectral shape but different absolute intensities. This may be due to the rough surface of the fiber affecting the contact of the AFM tip on the fiber surface for each spot.

As shown in FIG. 10, spectra C1, E1 (red spectrum) and C1', E1'(black spectrum) were plotted in the same figure. For comparison purposes, the intensity of the maximum peaks of all spectra was normalized to 1.0. By comparing C1 and C1', both spectra have ^{characteristic peaks at 1728 cm -1} and 1740 cm ^-1 , but the 1740 cm ^-1 peak at C1 has much higher intensity than that of C1'. Is observed, which indicates that fibers with a diameter of 263 nm have a higher β crystal content than fibers with a diameter of 390 nm throughout the core region. By comparing two color-coded pairs of curves, namely C1, E1 (solid line spectrum) and C1', E1'(dashed line spectrum), E1 and E1'compared to C1 and C1'. You will notice that it tends to always have a much higher 1740 cm- ^{1 peak.} This finding indicates that the morphology of the polymer chains is indeed inhomogeneous throughout the single fiber and that the shell contains more β-type crystals than the core. Interestingly, when comparing E1 and E1', there was no significant difference between the two spectra in terms of band shape, suggesting that the crystal structure or chain morphology in the shell is independent of fiber size. There is.

Further support for the conclusion that the rotating disc aligned electrospun PHBHx nanofibers have a core-shell structure can be obtained by examining the cross section of the fibers. 11a and 11a'show AFM height images of cross sections of two rotating disc aligned fibers from the same batch with diameters of 260 nm and 312 nm, respectively. Again, IR images were captured with two characteristic frequency of 1740 cm ^-1 and 1728 cm ^-1. In order to eliminate the effects of sample thickness variation and thermal drift, a ratio of the two mappings, I1740 / I1728, was taken as shown in FIGS. 11b and 11b'. Color bars show an increase in the relative intensity of IR absorption from purple to red. 2 Ton epoxy is because it has absorption IR laser negligible 1740 cm ^-1 and 1728 cm ^-1, IR mapping ratio in the epoxy region is should be equal to 1, which is the green / yellow in Figure 11b and Figure 11b ' expressed. As shown, in each of the two fiber cross sections, a clear red ring was observed at the periphery of the fiber cross section, resulting in a clear difference between the fiber cross section and the epoxy. This is because, in the ring region, greater than one ratio I1740 / I1728, or IR absorption at 1740 cm ^-1 indicating that higher IR absorption at 1728 cm ^-1. This finding confirms the presence of a thin shell containing more β-type crystals in the rotating disc aligned electrospun PHBHx nanofibers. In addition, it has been consistently found that the thickness of the ring is about 10 nm in both cross sections, which is comparable to the width of the red line along the fiber edge in FIGS. 8b and 8b'. This strongly suggests that the shell thickness in the fiber is about 10 nm independent of the fiber diameter. Further, in FIG. 11b, it is also observed that the color in the core region of the fiber cross section is yellow, which is shifted in the red direction as compared with the green color in the epoxy region. This indicates a slightly higher β-type content than the α-type in the core.

On the contrary, in FIG. 11b', the color in the core region is green, which is shifted toward purple as compared with yellow in the epoxy region, and shows β type having a lower content than α type. This finding, 1740 cm ^-1 peak in the spectrum C1 is consistent with the findings in Figure 10 is higher than 1740 cm ^-1 peak at C1 '. However, the ring is not perfect along the periphery of the fiber cross section (FIG. 11b) and there are some extra red regions beyond the ring (FIG. 11b'). This is probably due to the rubbing of the polymer during the microtome process, which results in random bursting or shrinkage of the polymer if the diamond knife does not cut well or the cutting speed is slow.

Based on experimental observations, we do not want to be bound by any particular theory, but the possible mechanisms for the formation of polymorphic inhomogeneous core-shell structures in electrospun PHBHx nanofibers collected on rotating disks are herein. Proposed in, and shown in FIG. The randomly coiled dissolved polymer chains are highly pulled during electrospinning, resulting in chain elongation and orientation along the tensile direction. As a result, the amorphous fibers that reach the rotating disc consist of oriented polymer chains along the fiber axis, which are plasticized by the residual solvent. During collection, the amorphous fibers are further pulled by the additional stretching force provided by the high speed rotating wheel. At the highest elongation, most of the polymer chains in the amorphous fiber are fully elongated and have a planar zigzag structure as shown by the fiber immediately after spinning in FIG. Due to the extremely rapid solvent evaporation rate on the fiber surface, the planar zigzag chains near the surface are kinically fixed and crystallize into a metastable β-type. Therefore, a thin layer of coating or shell is formed. The solidified shell then acts as a semi-permeable barrier, limiting evaporation of residual solvent in the core. As a result, especially some of the chains in the core are converted to a more stable spiral structure (alpha type). Therefore, the final solidified fiber has a core-shell structure, and the crystal structure is different between the core and the shell.

A core-shell model of the spatial distribution of α-type and β-type polymorphs can greatly facilitate the understanding of the structural / processed / characteristic relationships of PHBHx nanofibers. Α-type and β-type crystal structures having the same chemical composition but different molecular packings have been reported to have unique properties including mechanical properties, biodegradability, and piezoelectricity. For example, it is widely recognized that β-type P (3HB) has much higher strength and modulus than its α-type equivalent. Studies on the enzymatic degradation of P (3HB) revealed that the rate of degradation of the β phase with an all-trans structure is higher than the rate of degradation of the α phase with a spiral structure. More interestingly, recent experimental results indicate that the piezoelectric response of PHBHx nanofibers probably correlates with the introduction of β-type crystal structures. Therefore, the final properties of PHBHx nanofibers can be greatly influenced by the composition of the α and β phases, which are highly dependent on the electrospinning conditions. According to the core-shell model, if the application of PHBHx nanofibers requires properties dominated by β crystal structure, the radius can be increased by increasing the overall tensile force during electrospinning or by promoting solvent evaporation. The absolute content of β-type can be increased by making it faster than solvent diffusion in the direction. Furthermore, in order to increase the relative volume fraction of the β-phase-rich shell, the β-type relative content can be increased by reducing the fiber diameter.

Conclusion The spatial distribution of crystalline polymorphs in single electrospun nanofibers was studied for the first time using AFM-IR technology. 2 ₁ thermodynamically stable α-type consisting of a chain having a helical structure, and were investigated electrospun PHBHx nanofibers comprising two crystalline polymorphs of consisting of chains having a planar zigzag structure metastable β-type. Imaging of AFM-IR spectra and single PHBHx nanofibers demonstrated coexistence of α-type and β-type polymorphs on a single fiber scale, reconfirmed by SAED results. Furthermore, it was confirmed that the molecular orientation level and concentration of β-type are highly dependent on the fiber diameter. More importantly, AFMIR spectra and imaging revealed that the two polymorphs are spatially distributed as an inhomogeneous core-shell structure consisting of an α-type rich core and a β-type rich shell. .. Shell thickness remained constant throughout fiber size variability, indicating that shell formation is primarily controlled by competition between solvent evaporation and diffusion. Based on the above experimental observations, a possible mechanism of formation of the core-shell structure has been proposed. During fiber solidification, the planar zigzag chains resulting from the highly oriented free chains in the fiber were kinically fixed near the fiber surface, forming a beta-rich shell due to the extremely high solvent evaporation rate on the surface. .. The zigzag chains in the core region of the fiber are relaxed and converted into more stable spiral chains to form alpha-rich cores. This was possible due to the presence of residual solvent, the evaporation of which was inhibited by the tightly packed shell. In this study, AFM-IR technology is certainly an effective and efficient tool for nanoscale exploration of single electrospun fibers, with spatial resolution well below the diffraction limit in infrared for topographical and structural information. Shown to provide both. This study was considered as a template for nanoscale structural studies of various polymorphic materials where electrospinning promotes the formation of metastable crystalline phases, such as nylon-6 and polyvinylidene fluoride (PVDF). can do. Investigation of polymorphic distribution of nanofibers according to processing / collection conditions provides a deeper understanding of molecular chain behavior under extremely high shear / tensile forces and extremely rapid solvent evaporation during electrospinning. This level of basic understanding of structural / property / process relationships is an important first step towards the rational design and manufacture of polymer nanofibers with specific properties depending on the intended use.

Measurement of Piezoelectricity on Electrospun Poly [(R) -3-Hydroxybutyrate-co- (R) -3-Hydroxyhexanoate] (PHBHx) Nanofibers Using an electrospun high speed rotating disk as a collector , A macroscopically aligned PHBHx nanofibers containing a semi-stable β-type crystal structure were produced. Details of the electrospinning process can be found in Example 1. Specifically, the rotating disc was carefully wrapped in non-adhesive aluminum foil so that the self-supporting ribbon of highly aligned PHBHx nanofibers could be easily detached from the flat edges of the rotating disc.

It has been reported that the metastable β-type crystal structure of annealed PHBHx can be reannealed to α-type by heating to 130 ° C. Therefore, to facilitate the conversion of any β-polymorph to a thermodynamically stable α-polymorph, the ribbon of fibers was annealed in the oven at 130 ° C. for 24 hours. To maintain the macroscopic alignment of the fibers, the two ends of the fiber ribbon were clamped to a glass slide and held there during annealing.

Textile Matte Characteristic The morphology and crystal structure of the rotating disc aligned fibers before and after annealing were characterized by scanning electron microscopy (SEM) and wide-angle X-ray diffraction (WAXD).

Measurement of Piezoelectric Response A piezoelectric cantilever was used to test the piezoelectric response of macroscopically aligned PHBHx nanofibers. As shown in FIG. 13 (a), two piezoelectric lead zirconate titanate (PZT) sheets (T105-H4E-602, Piezo Systems) A and B are made of rectangular stainless steel having a length of 32 mm and a width of 3.5 mm. Adhered to the top and bottom of the steel layer. The top PZT sheet A has dimensions of 22 mm (length) x 3.5 mm (width), and the bottom sheet B is slightly shorter, measuring 12 mm x 3.5 mm. A thin brass cylinder with a diameter of 1 mm was glued to the tip of the cantilever to function as a test probe. Further details regarding the piezoelectric cantilever can be found elsewhere. Due to the inverse piezoelectric effect, the application of an external voltage across the PZT sheet A results in an axial displacement of the cantilever and test probe. FIG. 13 (b) shows the displacement of the probe at various applied voltages. The probe displacement was measured using a laser displacement sensor (LC-2450, Keyness Corporation) with a resolution of 0.5 μm. The external voltage was supplied by a function generator (Agilent 33220A, 20 MHz).

To measure the piezoelectric response of the fibers, a ribbon of fibers 5 mm wide x 10 mm long was bonded to a flexible PDMS substrate. The PDMS substrate was coated with insulating Kapton® tape to reduce background noise. Conductive copper tape (3M 3313, 1/2 inch) acting as an electrode was used to glue the two ends of the ribbon of fibers onto the substrate so that the aligned fibers were perfectly straight without load. .. The piezoelectric test was initiated by carefully engaging the cantilever probe onto the surface of the fiber ribbon and applying a square wave voltage supplied by a function generator across the cantilever to generate cantilever vibrations at a frequency of 10 Hz. .. Therefore, the probe in contact with the ribbon of the fiber tuned in and deformed the straight fiber. The induced voltage generated by the deformation of the aligned fibers was recorded on a digital oscilloscope (Agilent Infinitum 1.5 GHz, 8 GSa / s). The experimental configuration is shown in FIG. Control experiments were performed to verify that no apparent voltage signal was observed without the ribbon of fibers. All experiments were performed on a Newport optical table (RS100, Newport Corporation) to minimize background vibration.

Results and Discussion The morphology and crystal structure of the macroscopically aligned PHBHx nanofiber bundles before and after annealing are shown in FIG. After annealing, it is observed that the macroscopic alignment of the fibers remained constant. As evidenced by the loss of the β-type X-ray diffraction peak at 19.6 ° shown in FIG. 3b', the β-type disappeared, probably due to the thermodynamically stable conversion to the α-type. Furthermore, after annealing, the two diffraction peaks of type α became narrower, that is, the FWHM of the two α peaks became smaller, because the size of the α crystal was due to the reorganization of the molecular chains during annealing. It shows that it increased due to it.

The piezoelectric response of PHBHx nanofibers before and after annealing was measured at an applied voltage of 10 V. The results are shown in FIG. Quantitatively, the pre-annealed fibers (black signal) show a much higher voltage output (about 240 mV between peaks) than the post-annealed fibers (red signal), which is the piezoelectricity and β-type of the fibers. It shows a strong correlation with the presence of crystal structures.

The β crystal morphology contains a molecular chain exhibiting a planar zigzag structure. As shown in FIG. 5b, an electronegative O atom from the C = O group and an electropositive _{H atom from two} adjacent CH groups are present on the opposite side of the polymer backbone. As a result, a net dipole moment is formed perpendicular to the polymer backbone. Previous studies have shown that polymer chains in rotating disc aligned fibers are highly oriented along the fiber axis. Therefore, when the aligned fibers bend, the oriented polymer chains trapped in the fibers deform, resulting in a change in the dipole moment. As a result, a potential difference is created between the two ends of the fiber, which appears as a positive black voltage spike in FIG. On the other hand, the electric dipole is present in the α-type molecular chain of left-handed 2 ₁ helical structure aligns to cancel all dipole moments from each other (Fig. 5a). This results in a net dipole moment of zero in the α crystal, so annealed aligned fibers containing only the α-type crystal structure do not exhibit a piezoelectric response. Interestingly, Ando et al. (18-20) investigated the piezoelectric properties of α-PHBs and found that oriented α-type PHBs exhibit intrinsic shear piezoelectricity. _{The internal rotation of the dipole atomic group (O = C-CH 2} , FIG. 5a) associated with the asymmetric carbon atom in the α-type spiral chain appears to result in electric polarization when shear stress is applied. However, this shear piezoelectricity was not observed in FIG. This difference may be due to: (1) the molecular chain is not sufficiently sheared to exhibit shear piezoelectricity in this experiment due to the low shear piezo coefficient of α-PHB; or (2) the annealing process It is possible that the high alignment level of the chain was reduced, so the annealed α-crystals are randomly distributed in the fiber. Further investigation is needed to gain better insight into this finding.

Another interesting feature in FIG. 4 is the alternating appearance of positive and negative voltage spikes. This phenomenon can be explained by the charging and discharging of piezoelectric PHBHx fibers during the press-hold-release cycle (FIG. 6). As the brass probe begins to press the ribbon of the fiber, the oriented molecular chains within the fiber deform, which results in a momentary potential difference between the two ends of the ribbon of the fiber. Quantitatively, this potential difference is represented in FIG. 6 by a positive voltage signal of 119 millivolts. In response to this induced potential difference, an external free charge (green line) is driven by the ribbon of the fiber to cancel this potential difference. During this process, the net charge (pink line) on the ribbon of the fiber increases. When the press is held, i.e., the mechanical strain is kept constant, the piezoelectric constrained charge is offset by the external free charge, so that the induced potential difference and the net charge gradually decrease. With zero potential difference, the ribbon of fibers is fully charged. When the press is released, the incorporated potential difference disappears and the external free charges accumulated at both ends of the ribbon of the fiber flow in the opposite direction of the accumulation process. This reversal results in a quantitative opposite potential difference (-123 millivolts), which is observed as a negative voltage signal. As shown in FIG. 6, the opposite potential difference and net charge are also gradually reduced. As soon as the potential difference reaches zero, the fiber ribbon discharge is complete.

The above experimental results suggest the possibility of using piezoelectric PHBHx nanofibers as nanogenerators. On the other hand, the piezoelectric properties make the fiber ribbon a promising component of the piezoelectric sensor. To assess the sensitivity of the ribbon of the fiber, the applied voltage is adjusted to explore the relationship between the induced voltage and the applied pressure using the device shown in FIG. The pressure applied on the fiber is calculated by the following formula.
p = (K × D) / A
In the equation, K represents the spring constant of the cantilever, D is the axial displacement of the probe, and A is the contact area between the probe and the ribbon of the fiber. It is known that K = 128.9 ± 2.2 N / m22.

D is measured by a laser displacement sensor as shown in FIG. 1b, where A = πr ² and r = 0.5 mm. The induced voltage is recorded in the applied voltage range of 4-10V and the increment is 1V. A plot of the induced voltage as a function of the applied voltage is shown in FIG. 7b, where a linear relationship was observed. Therefore, the sensitivity of the ribbon of the fiber can be calculated from the inclination of the linear fitting, and it can be seen that it is 7.46 mV / kPa. This sensitivity seems to be much lower compared to other piezoelectric sensors, but this is 1) changing the configuration of the sensor, eg stacking multiple layers of piezoelectric fiber ribbons on top of each other; or fiber. It can be improved by a method comprising increasing the concentration of the piezoelectric active component (β crystal) in.

Conclusion In this study, we investigated the piezoelectric properties of electrospun PHBHx nanofibers according to their crystal structure. The piezoelectric response of the rotating disc aligned fibers before and after annealing was characterized. The results showed that nanofibers containing a metastable β-type crystal structure consisting of planar zigzag chains show a clear piezoelectric response (240 millivolts between peaks) when mechanically deformed. However, after annealing and conversion of β-type crystals to α-polymorphs, this piezoelectric response disappeared. This finding showed a strong correlation between the piezoelectric properties of the fibers and the presence of β-type crystal structures. Then, the sensitivity of the piezoelectric PHBHx nanofibers to the applied pressure was measured. When the induced voltage was recorded when the applied voltage was changed, a linear relationship was observed. From the inclination, the piezoelectric sensitivity of the fiber was measured as 7.46 mV / kPa. The implications of these preliminary studies on the piezoelectricity of PHBHx are broad. The piezoelectric performance of this material can be significantly improved by increasing the concentration of the active β-type crystal structure. That goal can be achieved by utilizing innovative polymer processing techniques such as those discussed herein. Piezoelectric PHBHx outperforms all other piezoelectric polymers in terms of excellent biodegradability and biocompatibility, environmental friendliness, and most importantly, low manufacturing costs. It is a highly promising piezoelectric polymer that can be found in many advanced fields, including portable and foldable electronic devices, artificial electronic skin, and implantable sensors.

Stress-Inducing β-Crystating Material of Poly [(R) -3-Hydroxybutyrate-co- (R) -3-Hydroxyhexanoate] in Membrane Poly [(R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] poly [(R) 3.9 mol% and 13 mol% 3HHx copolymer content. The (R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate] (PHBHx) copolymer was supplied by Procter & Gamble Company, Cincinnati, Ohio and was not further purified. The weight average molecular weights of the 3.9 mol% and 13 mol% 3HHx copolymers were 843 kg / mol and 792 kg / mol, respectively. Chloroform was added to Sigma-Aldrich Co., Ltd. , Purchased from Ltd and used without further purification.

Sample Preparation A PHBHx membrane was prepared by dissolving the polymer in chloroform and solvent casting the membrane onto a Teflon® block to avoid binding to the cast surface. The annealed membrane was adjusted in an oven at 70 ° C. for 4 hours before analysis. Samples for IR spectroscopy were cast from 3 wt% PHBHx solution and samples for Raman and XRD analysis were cast from 10 wt% PHBHx solution. IR measurements required lower concentrations to produce a film thin enough to avoid total absorption of the IR beam. Conversely, Raman and XRD needed higher concentrations to provide a sample thick enough for Raman to focus the laser and XRD to generate diffraction. The cast membrane was dried in a vacuum chamber for about 4 hours to remove all residual chloroform. After removal of the solvent, the membrane was melted in a convection oven at 140 ° C. for 20 minutes to completely melt the polymer and then immediately quenched in ice water to maintain an amorphous state below Tg. Membranes were stored at -25 ° C, loaded into a mechanical tensioner at -25 ° C, pulled to about 50% strain at -25 ° C and returned to room temperature. IR or Raman spectra were collected with the final strain applied continuously up to about 150% strain. After the sample was completely distorted, XRD diffraction profiles were collected for the sample and measured by Raman.

Fourier Transform Infrared (FTIR) Spectroscopy Infrared absorption spectroscopy using a Thermo Nicolet 670 Nexus FTIR spectrometer equipped with a DTGS KBr detector and a KBr beam splitter that transmits a polymer film mounted on a mechanical tensile device. Was recorded. Spectra were collected by simultaneously adding 16 scans at a resolution of ^{4 cm -1} over a range of 4000~600cm ^-1. IR spectra were processed by baseline correction in a tertiary spline fitting, the ^{spectra were truncated to 1550-700 cm -1} , and they were normalized using SNV normalization. The truncation was performed to avoid distortion of the normalization coefficient due to carbonyls that are excessively absorbed due to the thickness of the membrane.

Raman spectroscopy Raman measurements were made using a Raman instrument consisting of a diode laser (Ondax) with excitation at 785 nm and an Ondax probe head optical filtering module. The collected scattered light was analyzed using a Kaier Optical Systems Hollowspec 1.4 spectroscope equipped with an Andor CCD system. Raman spectra were processed using Kasier Optical Systems Holopec software. The spectra were collected with an exposure time of 20 seconds and 10 exposure integrations and subjected to equivalent dark subtraction. Raman spectra were processed by dividing by the spectrum of a white light source, baseline corrected using cubic spline fitting, and normalized using SNV normalization.

Wide-angle X-ray diffraction (WAXD)
Bruker D8 XRD with a Cu tube source was used at 40 kV and 40 mA to generate a 1.5418 Å X-ray beam and the wide-angle X-ray diffraction profile was measured. The LYNXEYE_XE detector was used in OD mode to detect scattered X-rays. WAXD profiles were recorded in a 2θ range of 10-25 ° or 10-30 ° in 2θ increments of 0.025 and collected 1 second at each point. Data acquisition was cycled and simultaneously added until a reasonable signal vs. noise level was achieved. The polymer film was measured with a mechanical tensioner with the X-ray beam perpendicular to the strain direction.

Results and Discussion β-type characterization in PHA is a somewhat controversial area. The β-type, which is typically discussed as the crystalline form of the polymer, is not a true crystalline form, but rather a metastable partial filling of polymer chains having the same structure. Therefore, the structure of the unit cell is unknown, and the lattice constant is not defined except for c (fiber cycle) = 0.470 nm. Conversely, the α crystal morphology shows an α (020) peak at 2θ = 13.3 °, an α (110) peak at 2θ = 16.7 °, and an α (101) peak at 2θ = 21.1 °. The d-spacing of α-peaks results in overlap between α-type and β-type when tested with XRD. This duplication sometimes obscures the diffraction test to confirm the presence of β-type, especially in the case of one-dimensional WAXD. Due to the difference in polymer chain structure between α-type and β-type, IR and Raman spectroscopy should be used as comparative analysis methods, and diffraction tests are for confirming the presence of β-type. In the skeletal region of the spectrum, there are several vibrational modes due to COC and CO expansion and contraction in the polymer chain, which have varying vibrational frequencies between the two phases. However, due to the simplicity of the single band test, it may be preferable to observe the carbonyl peak to determine the β content, as performed for the α vs. amorphous content of the polymer. The problem is that the location of the β carbonyl peak is still controversial. One group has reported the oscillation frequency of about 1730 cm ^-1 or groups have reported the vibration frequency of 1740 cm ^-1. The second peak position is the most problematic, because any analysis of the β content of the PHA sample using the carbonyl band, if the carbonyl oscillation mode in β type is really the same as in the case of amorphous polymers, This is because it can be similarly correctly interpreted as vibration in the amorphous content of the polymer. For analysis using Raman, IR, and WAXD for the purpose of differentiating the position of the β-type carbonyl band, a film of 13 mol% 3HHx PHBHx with strain-induced (strain-induced) β-type was generated.

As mentioned earlier, XRD results can be obscured in determining the presence of type β, but it is still important to obtain a diffraction profile for comparison with type α. Moreover, in highly oriented samples, most of the α diffraction peaks tend not to be observed in one-dimensional diffraction tests due to the crystal orientation in the sample. FIG. 1 includes diffraction profiles of untreated 3.9 mol% and 13 mol% 3HHx PHBHx polymer membranes after annealing and 13 mol% 3HHx PHBHx polymer membranes after strain is applied. A 3.9 mol% 3HHx PHBHx annealing film was included in the test to show the number of α crystal diffraction peaks for comparison with the pulled film.

The 3.9 mol% 3HHx PHBHx WAXD profile contains another α peak of 2θ up to 19.5 in addition to the three diffraction peaks previously listed, while the 13 mol% 3HHx PHBHx annealing film is 21.5. It includes α (020) and (110) peaks in addition to a wide range of peaks centered on 2θ. The WAXD profile of the pulled membrane also contains α (020) and (110) peaks, but shows a larger and wider peak centered on 2θ at 19.5. This wide diffraction peak is typical for PHA samples containing β-type. However, it should be noted that the higher 2θ diffraction peaks of the α crystal are almost unobservable. The loss of diffraction from other crystal planes is due to the high degree of orientation introduced into the sample during the two stages of tension. Since the sample is oriented in the tensile direction perpendicular to the X-ray beam in XRD, the α crystal lamella is oriented in the c-axis parallel to the tensile direction. Furthermore, the β type is also oriented in the c-axis parallel to the tensile direction. Comparisons with the other diffraction profiles in FIG. 1 show problems with the use of XRDs to confirm the presence of β-type. There is a diffraction peak originating from the α type at the same position as the maximum value of the β peak, which is found in the 3.9 mol% 3HHx PHBHx film. If the sample is oriented in a particular fashion with respect to the X-ray beam path, this peak will have increased relative intensity, especially if the diameter of the α crystal is small and this widens the diffraction peak, it will be of type β. May cause illusions. Overall, these profiles strongly suggest that β-type was produced in the pulled membrane, but further analysis was needed.

To confirm the presence of β-type in the 13 mol% 3HHx PHBHx pulled membrane, IR spectra were measured when the membrane was pulled at room temperature after the first strain was applied below Tg. These spectra only contain the skeletal region because the membrane was too thick and resulted in overabsorption of the carbonyl. Thinner samples are not mechanically stable enough for use in mechanical tension devices. FIG. 21 includes an IR spectrum of a net 13 mol% 3HHx PHBHx polymer membrane as a function of increasing strain.

All ^1278Cm -1 attributable to the spiral structure, 1263cm ^-1, and based on the presence and consistency of the band at 1228cm ^-1, polymer crystallized α-type before the spectrum acquisition begins. Chaturvedi et al. Compared the IR and FT-Raman spectra of PHB samples containing β-type with quantum chemical calculations of the vibrational dynamics of linear zigzag PHB chains to determine the vibrational band associated with β-type of the polymer. Their analysis yielded a number of vibrational frequencies, some of which were not a function of the structure of the polymer chain but contained bands in the IR spectrum that increased with increasing strain. Specifically, in the spectrum, the bands 1305 cm ^-1 , 1142 cm ^-1 , 1080 cm ^-1 , 969 cm ^-1 , and 911 cm ^-1 all increased as the membrane was pulled. All of these vibration modes were calculated for the β structure, not due to the amorphous or α crystalline phase of PHB. Along with the results from XRD, IR spectra support that the pulled PHBHx can induce β-type just as observed in homopolymers, despite the high 3HHx content of the sample.

XRD and IR analysis revealed that β-type can be formed on the PHBHx membrane when strain is applied, but carbonyl observation is not possible due to the limitations imposed on IR measurements by sample thickness. rice field. Carbonyls are heavily influenced by their local environment and are often used to determine the presence of amorphous and α crystalline phases in PHA. Specifically, the α-crystal IR band centered at 1720 cm ^-1, amorphous band centered at 1740 cm ^-1. For this reason, Raman spectra were also collected for the pulled membrane. Raman spectroscopy does not measure light absorption like IR spectroscopy, but it does measure the so-called Raman scattering of the material. When the light from the laser in the Raman spectrometer interacts with the sample, photons can bind to the vibrational mode of the functional groups of the sample. The energy difference between the photons before and after this bond corresponds to the energy difference between the two resonant states of the material. Plotting the intensity of scattered light as a function of the frequency difference between scattered and incident photons forms a Raman spectrum. Therefore, this method does not require light to pass through the sample, so strong peaks cannot be excessively absorbed due to the thickness of the sample. In addition, IR is much more sensitive to changes in dipole moments, while Raman is more sensitive to changes in polarizability, so in Raman the carbonyl stretching vibration mode is much weaker than IR. .. FIG. 22 contains Raman spectra of untreated 13 mol% 3HHx PHBHx polymer films before and after straining.

The spectra appear somewhat different compared to the spectra in FIG. 21 because the photons in the Raman and IR spectra interact with the sample differently. However, the vibration frequency of the band should not change between methods. With that in mind, the changes in the spectrum with increasing strain are very similar to those of IR. Specifically, 1451 cm ^-1 , 1379 cm ^-1 , 1354 cm ^-1 , 1177 cm ^-1 , 1080 cm ^-1 , 966 cm ^-1 , 908 cm ^-1 , 858 cm ^-1 , 627 cm ^-1 , 478 cm ^-1 , and 454 cm ^-1 . All bands increase when strain is applied to the polymer membrane. All of these vibration bands are attributed to the beta structure of the polymer and serve to reinforce the results of XRD and IR analysis. To better observe the carbonyl peak in Raman, FIG. 22 shows a magnified spectrum of untreated 13 mol% 3HHx PHBHx polymer membranes before and after strain, with a focus on the carbonyl region.

When strain is applied to the film, an exchange occurs between the α crystal peak and the amorphous peak, and it seems that another peak exists between them. The intensity of the α crystal band decreases, while this new band and amorphous peak increase. The position of the new peak is about 1730 cm ^-1 , which is the same position reported by Chaturvedi et al. For the β-type structure. These results, the correct position of the β-carbonyl peak is 1730 cm ^-1, suggesting strongly that it is not a 1740 cm ^-1. However, since the Raman scattering of the carbonyl functional group is weak, there is a large amount of noise in the spectrum, and it is difficult to decompose a new band. Further research is needed in this area to produce a better spectrum as a function of percent distortion. Finally, tools such as 2DCOS for such datasets, not only to determine the carbonyl band for β-type, but also to analyze the process by which β-type is formed when the polymer is pulled. A more thorough analysis using is to be applied.

Mechanical tensioning of membranes for the formation of β-type of poly ((R) -3-hydroxybutyrate-co- (R) -3-hydroxyhexanoate) (PHBHx) Experimental procedure 10 weight percent PHBHx in chloroform Solution was prepared. The Hx content of PHBHx was 13 mol percent. The membranes were then cast onto glass slides at room temperature using 2.5 mL of solution for each membrane. These membranes were dried at room temperature for at least 1 hour and then placed in a vacuum chamber overnight.

After overnight in a vacuum chamber, the membrane was placed on a Teflon block and melted in a furnace at 140 ° C. for 20 minutes. Immediately after removal from the furnace, the membrane was quenched in ice water and transferred to the freezer. These membranes were shown to be amorphous by ATR FTIR and XRD measurements.

The membrane was then removed from the freezer and left at room temperature for the desired time. The membrane was then pulled at room temperature as shown in FIG. 24 and then analyzed.

Results and Discussion Planar Zigzag Formation We succeeded in forming a planar zigzag structure of PHBHx by one-step tension at room temperature following isothermal crystallization at room temperature. The pulled membrane behaved consistently in a similar manner in forming the β structure. Most of the membrane did not try to stretch and remained unchanged. This section of the membrane does not result in the formation of β-structures. The section of membrane adjacent to the clamp of the tensioner stretched to form a distinctly different transparent area, a process called "necking". This pulled region is where the β structure is formed. First, the β region of the membrane tended to break at low elongations. To alleviate this, tape was attached to the ends of the membrane prior to mounting on the puller. This reduced the direct contact between the membrane and the clamp, allowing the membrane to be pulled further before breaking.

The β structure was confirmed using XRD and Raman spectra. In the XRD spectrum, the peaks at 2θ values of 13 and 16 correspond to the α crystal planes of 020 and 110, respectively. The apparent peak at the 2θ value of 22.5 is actually the overlap of the three peaks from the α crystal structure. The peak of 19.2 is caused by the presence of a planar zigzag structure ^7,8 . FIG. 25 shows the XRD spectra of the unstretched portion and the pulled portion of the same membrane, respectively. It can be clearly seen that the mechanical strain forms the β structure. Moreover, the size of the α crystal in the stretched region of the membrane is smaller than the size of the α crystal in the non-pulled region. In the stretched region of the membrane, the average α crystal domain size was 21.3 nm for the 020 plane and 20.4 nm for the 110 plane. All crystal domain sizes are calculated using Scherrer's equation with a K value of 1. This suggests that the formation of a planar zigzag structure also results in a reduction in the size of the α crystal. Such results have been seen in other PHA as well.

The size of the β crystals formed by this procedure is small, as can be seen by the wide peaks corresponding to the β structure. Quantitatively, the size of the β crystal domain is 3.8 nm on average. This raises the question of whether true β crystals are formed. Instead of a true crystal, this procedure is thought to result in the formation of a regular planar zigzag phase with too few aligned chains to form the crystal. This may be due to the loss of alignment in the amorphous prior to crystallization and the formation of α-crystallites.

Planar zigzag reversibility A membrane that is pulled at room temperature to form a planar zigzag structure and then released from mechanical stress does not maintain a planar zigzag structure. After the tension is released, only α crystals are maintained. However, after the tension is restored, the β structure returns. This effect can be seen in FIG. 27 for each of the pulled membrane, the released membrane, and the immediately re-tensioned membrane. In addition to the XRD spectrum, a Raman spectrum of the released membrane was obtained, and it was confirmed that the β structure disappeared.

This was not the case when the released membrane, which had been left at room temperature for 24 hours, was pulled again. As can be seen in FIG. 28, the β content gradually increased as the elongation increased. Due to the increased crystallization, when the membrane was pulled again, the necked area continued to grow as the membrane was stretched. The surrounding non-necking area was too crystalline and rigid to neck, so the already-necked area continued to grow. It can be seen that the size of the peaks associated with the α crystal structure decreases before the peaks associated with the β structure increase. The size of the α crystal domain also seems to decrease slightly, but this is not as pronounced as the overall decrease in peak height. This indicates that when the film is pulled again, the α crystals first melt into an intermediate state and then form a β structure.

It is clear that the β structure was not maintained after the tension was released.

Further annealing was performed at 48 ° C. for 2 hours. The XRD spectrum of this film is shown in FIG. The α content of the pulled membrane is significantly increased as a result of this treatment. In addition, the amount of β-structure decreased as a result of annealing. Once released, it seems that a small amount of β-structure can be maintained. However, the membrane is attached to the XRD sample holder and this process can apply some tension to the membrane and restore the β structure.

To optimize the formation of β-structure, the membranes were subjected to isothermal crystallization for different times and then pulled. The membrane was pulled after 25 minutes, 28 minutes, 30 minutes, 35 minutes and 40 minutes. When pulled after 25 minutes of crystallization, the entire membrane stretched over 200% with no sign of breakage. The pulled membranes after 25 and 28 minutes of crystallization were analyzed using XRD and Raman, as seen in FIGS. 30 and 31, respectively.

The above spectrum shows that the membrane pulled after 25 minutes of crystallization did not form any β structure, while the membrane pulled after 28 minutes formed a large amount of β structure.

In this example, it was shown that isothermal crystallization of 13 mol percent Hx in PHBHx membrane at room temperature could form a β structure using a single tension. However, the XRD peaks of this structure are wide and it is unclear whether this β structure is in the form of very small crystals or in the form of amorphous, regular phases.

Recording XRD and Raman spectra as a function of time while the membrane is annealed at various temperatures provides insight into these changes and how the β-structure is fixed after tension is released.

Further, the examples disclosed herein are for PHBHx with a Hx content of 13 mol percent. This is a large amount of Hx monomer and many compositions of PHBHx with lower Hx content are available. These polymers behave differently from 13 mol percent PHBHx, and parallel studies of each of these PHBHx compositions will yield interesting results regarding the relationship between Hx content and β-structure formation.

CONCLUSIONS: A novel method of mechanical tension following room temperature isothermal crystallization has shown that the β structure of PHBHx is formed in PHBHx films with an Hx content of 13 mol%. After 28 minutes of crystallization, different tensile regions of the membrane were seen, which caused a large section of the membrane to form a β structure. Earlyer tensions do not result in any β formation, and later tensions do not form equally wide sections of β structure, resulting in a rigid film that is prone to breakage. XRD and Raman analysis confirmed that α crystals must be present prior to tension in order to form a β structure, and that the size of α crystals is reduced during this process. This is consistent with the formation mechanism suggested by previous literature on α crystal tie points.

The β structure has been shown to be reversible and the tensile process differs between the initial β formation and re-tensile processes. In the initial formation process, β forms immediately and does not change with elongation. Rather, more of the membrane is converted to β-structure as elongation increases. After being released, the β structure disappears. When the membrane is left at room temperature for 24 hours and then pulled again, the β structure gradually grows as it stretches, due to the fact that the rest of the membrane is rigid and only the previously pulled section can stretch. Form to. However, in the stretched section of the membrane, the α content does not increase dramatically, while in the non-pulled section of the membrane, the α content increases dramatically.

The β structure could also be annealed to the α structure again at a temperature as low as 48 ° C, but interestingly the size of the α crystal did not increase. Therefore, the number of α crystals should increase during this process without changing the size.

Since the β structure of PHBHx is easily formed in the membrane under easily achievable conditions, it is expected to develop into a commercial product in the future. Further investigation may realize the potential of this biodegradable, biocompatible, durable, and piezoelectric material for new products.

Poly in the gel film ((R)-3-hydroxybutyrate-co-(R)-3-hydroxyhexanoate) Tensile CHCl ₃ / DMF or CHCl _3/1 for the formation of β-type (PHBHx), A thermoreversible sol-gel transition was observed at PHBHx in 4-dioxane solution. PHBHx was _{dissolved in a two-component solvent system of CHCl 3} / DMF or CHCl ₃ / 1,4-dioxane at high temperature (100 ° C.) with stirring, while CHCl ₃ was recognized as a good solvent for PHBHx, while DMF and 1, 4-Dioxane is a poor solvent for PHBHx at room temperature. The clear PHBHx solution was gradually cooled to room temperature (about 20 ° C.) to produce a milky white gel. As can be seen in FIG. 32, increasing the temperature from room temperature to 100 ° C. resulted in a sol-gel transition, which was found to be reversible.

A thin film of the thermoreversible gel was formed by rubbing the thermoreversible gel on a glass slide and then drying the moist rubbed gel under ambient conditions. After evaporation of the solvent, a clear and smooth gel film was obtained. Again, the crystal structure of the gel film was characterized by WAXD and the resulting diffraction profile was found to be very different from that of the raw PHBHx powder and freeze-dried gel plotted in FIG. 33. Note that in FIG. 34a, the α (111) peak (2θ = 22 °) disappeared and the α (110) peak became wider and weaker. On the other hand, a new diffraction peak at 2θ = 19.7 ° appeared. This new peak is very similar in peak position and peak width (FWHM) to the β-type diffraction peak observed in the WAXD profile of rotating disc aligned fibers. More interestingly, the WAXD profile of the pulled gel membrane (FIG. 34b) shows an increase in the relative intensity of this new peak and the reappearance of the α (110) peak. These findings indicate that β-type crystallites may be present in the gel film and that tension in the gel film facilitates the development of both α and β crystal structures. Recrystallization of kinetic fixed molecular chains in the gel membrane during tension results in a change in enthalpy within the system, which causes the large amount of physically observed heat when the gel membrane is pulled. Can be explained.

These preliminary results indicate the appearance of β crystal structures within the gel membrane, which indicates that it may be possible to produce large quantities of PHBHx thin films with high concentrations of β crystals. These β-type rich PHBHx thin films could find many new uses, which were not the β-type crystallinity observed in α-PHBHx.

Although preferred embodiments of the invention have been shown and described herein, it will be appreciated that such embodiments are provided for purposes of illustration only. One of ordinary skill in the art will come up with a number of modifications, modifications and replacements without departing from the spirit of the invention. Accordingly, the appended claims are intended to include all such modifications as being within the spirit and scope of the invention.

Method for Making Polarized PHA Copolymer-based Articles 5 parts of PHBHx3.9 powder is _{dissolved in 95 parts of chloroform (CHCl 3} ) at 100-125 ° C., the solution is transferred to a tray and placed in a vacuum furnace. ^{The furnace is maintained at 10-3} torr and within the temperature range of 120-140 ° C. until a PHBHx solution with about 80% by weight PHBHx3.9 and 20% by weight DMF is obtained. The PHBHx3.9 solution is transferred as a membrane to a press, subjected to a pressure of 3000 PSI and heated to various temperatures in the range 95-125 ° C. The membrane is then rapidly cooled in an ice bath. The film is then transferred to a polaritizer consisting of two polished copper plates and the copper plate is connected to a high voltage DC power supply. The temperature of the membrane is raised slightly above the melting point to polarize the membrane. In one embodiment, during polarization, the temperature is linearly reduced at 2 ° C./min to 30 ° C./min and the polarizing electric field I is linearly increased from 25 KV / cm to 1000 KV / cm. At room temperature, reduce the polarizing electric field to zero.

Claims (21)

A method for preparing a polarized polymer composition, which is a method for preparing a polarized polymer composition.
a) A step of providing a layer of a polarizable polymer composition, wherein the polarizable polymer composition comprises a polyhydroxyalkanoate-based copolymer and a step.
b) The process of directionally perturbing the layer to induce polarization,
c) Optionally, a step of polarizing the polymer composition of the directionally perturbed layer by applying a high electric field with a strength lower than the strength that results in substantial dielectric breakdown of one or more polymers.
d) Optionally, a step of annealing the polarized polymer composition of the layer at a temperature lower than the melting temperature of the crystals of the polarized polymer composition, whereby the polarization is maintained to the crystal melting point of the polar crystals of the polymer composition. Process and
The method comprising the β-type PHA-based copolymer, wherein the polarized polymer composition is present in an amount of about 10% to about 99% as measured by X-ray diffraction. The step of providing the layer comprises electrospinning a ribbon of fibers from a solution of the polyhydroxyalkanoate-based copolymer in one or more solvents, each fiber of the ribbon of electrospun fibers in β-type. The method of claim 1, comprising a formed shell and an α-shaped core. The method of claim 1, wherein the step of providing the layer comprises forming a layer from a solution of the polyhydroxyalkanoate-based copolymer in one or more solvents. The method of claim 1, wherein the step of providing the layer comprises forming a layer from a melt composition of a polyhydroxyalkanoate-based copolymer. The method of claim 4, wherein the step of directionally perturbing the layer comprises calendering, shearing or cold stretching the layer after quenching. The method of claim 1, wherein the step of providing the layer comprises forming a film from the gel composition of the polyhydroxyalkanoate-based copolymer. The method of claim 6, wherein the step of directionally perturbing the membrane comprises drying the gel under shear pressure or freezing the gel to induce shear by crystallization of the solvent. Polyhydroxyalkanoate-based copolymers are hydroxybutyrate units, hydroxyhexanoate units, vinyl units, vinylidene units, ethylene units, acrylate units, methacrylate units, nylon units, carbonate units, acrylonitrile units, cellulose units, pendant fluoro, chloro. , Amides, acrylates and esters other than methacrylate units, cyanides, nitriles other than acrylonitrile units, or units with ether groups, protein units, and at least one monomer unit selected from the group consisting of combinations thereof. , The method according to claim 1. The configurable polymer compositions are polyvinyl chloride, polymethylacrylate, polymethylmethacrylate, poly (vinylidene cyanide / vinyl acetate) copolymer, vinylidene cyanide / vinyl benzoate copolymer, vinylidene cyanide / isobutylene copolymer, vinylidene cyanide. / Methyl methacrylate copolymer, vinylidene fluoride copolymer, polyvinyl chloride, polyacrylonitrile, polycarbonate, cellulose, protein, synthetic polyester and ether of cellulose, poly (γ-methyl-L-glutamate), vinylidene copolymer, nylon-3, nylon- 5. The method of claim 1, further comprising one or more polymers selected from the group consisting of 5, nylon-7, nylon-9, nylon-11 and blends thereof. Claimed that the step of polarization the polymer composition of an optionally directionally perturbed layer is carried out at a temperature of about 20 ° C. to about 120 ° C. for up to about 5 hours using an electric field of at least 1 MV / cm. The method according to 1. The method of claim 1, wherein the layer is annealed at a temperature in the range of about 125 ° C to about 150 ° C for at least 1 hour. A device comprising a PHA-based copolymer layer comprising at least one of a ribbon of electrospun fibers of a polyhydroxyalkanoate-based copolymer obtained by the method of claim 1 or a polarized polymer composition, wherein the layer is piezoelectric. Each of the ribbon and polarized polymer compositions of electrospun fibers, configured to show one or more of the effects, pyroelectric and piezoelectric effects, is about 10% to about 10% as measured by X-ray diffraction. Devices containing β-type PHA-based polymers present in 99% amounts. 12. The device of claim 12, which is a sensor configured to generate a potential difference or voltage in response to changes in the dimensions of the PHA-based copolymer layer. 13. A change in dimension is caused by one or more of the following properties on the surface of the PHA-based copolymer layer: humidity, temperature, salt content, nutrient attachment or injection, and metalloid attachment: Described device. The sensor comprises multiple sensor surfaces and / or interfaces so that each surface / interface monitors one of the following characteristics: humidity, temperature, salt content, nutrient attachment or injection, and metalloid attachment: The device according to claim 12, which is independently configured in. 12. The device of claim 12, an actuator configured to expand or contract upon application of charge across the PHA-based copolymer layer. A nanomotor comprising one or more piezoelectric actuators of claim 16. It is configured to be placed or implanted near the vasculature components of an animal or human patient, and the heartbeat of the animal or human patient produces dimensional changes in the PHA-based copolymer layer to operate the monitoring or treatment device. 16. The device comprising one or more actuators configured to generate a potential difference or voltage for the purpose. 18. The device of claim 18, wherein the treatment or monitoring device comprises a pacemaker, an insulin pump, an in-situ glucose monitor, or a blood pressure monitor. 18. The device of claim 18, wherein the vasculature component comprises a vein or artery that beats in a rhythm corresponding to the patient's heartbeat. The following properties: A step of providing a PHA-based copolymer sensor layer configured to monitor one of humidity, temperature, salt content, nutrient adhesion or injection, and semi-metal adhesion, the PHA-based copolymer. The sensor layer comprises at least one of a ribbon or polarized polymer composition of electrospun fibers of the polyhydroxyalkanoate-based copolymer obtained by the method of claim 1.
A step of exposing the PHA-based copolymer sensor layer to one or more of the properties such that the PHA-based copolymer sensor layer causes a dimensional change caused by one or more of the properties.
A method comprising the step of detecting a potential difference or voltage according to a change in the dimensions of a PHA-based copolymer sensor layer.

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